E&fSo/iSoi - A5S1 IM2-*P-Ml&h
ATOMIC ENERGY AUTHORITY, EGYPT
(Proceedings oft fie Jirst InternationalSpring School &Symposium
Editors Prof Dr. Louis N. Shehata Prof Dr. Abdel Wahab A. El-Sayed SAMS 94
VOL.11 CONTRIBUTED PAPERS 15-20 MARCH1994 CAIRO, EGYPT We regret that some of the pages in the microfiche copy of this report may not be up to the proper legibility standards, even though the best possible copy was used for preparing the master fiche
FIRST INTERNATIONAL SCHOOL & SYMPOSIUM ON ADVANCES IN MATERIALS SCIENCE (SAMS 94)
UNDER THE AUSPICES OF HIS EXCELLENCY ENGINEER
MM. mAZA
MINISTER OF ELECTRICITY AND ENERGY JJedicaticn
The Materials Science and. Engineering Community in Egypt is presenting this symposium in the Honour of
2Pu>£. ^.31. Jlammad Chairman. Atomic Energy Authority On his 60th Birthday Anniversary
In recognition of his outstanding record and his sustained effort in promoting educational, RSiD and industrial activities in the field of materials science and engineering at the Egyptian Atomic Energy Authority as well as universities, research institutions and industry. Prof. Hammad contributed to the education of fundamentals of materials science in several Egyptian univerisites and his students ranged from physicists and chemists to metallurgical, chemical, nuclear, mechanical and electrical engineers. A good number of these students are holding prominent academic and industrial posts in Egypt and on the international level. Several RSiD activities and programmes in various areas of materials science and engineering >> the AEA and elsewhere in Egypt haw been of his initiation and/or enjoying a great share of his contribi.'.ion. This has been published in national and international journals in more than one hundred publications. Prof. Hammad have always preached for and practically contributed to the interaction between academic institutions and industry and also to promoting approaches to develop materials strategies in Egypt.
i
SAMS 94
CHAIRMAN: PROFESSOR F.H. HAMMAD CO-CHAIRMAN: PROFESSOR H.F. ALY
ORGANIZING COMMITTEE
PROF. A.A. EL-SAYED PROF. LOUIS N. SHEHATA PROF. I.D. ABDEL RAZEK PROF. M. KHORSHID FAYEK PROF. S.M. EL-RAGHY PROF. H.I. SHABAN PROF. M. EL-SAYED ALI
- iii -
SCIENTIFIC COMMITTEE
CHAIRWOMAN PROF. V.K. GOUDA MINISTER OF SCIENTIFIC RESEARCH 1. INTERNATIONAL
PROF. V.L. AKSENOV (RUSSIA)
DR. CN. DE CASTRO (PORTUGAL)
DR. DEDINKO (RUSSIA)
PROF. A. FAHMY (USA)
DR. H.P. FRITZER (AUSTRIA)
PROF. P. HREN (USA)
DR. D. JENSEN (DENMARK)
DR. BUCHKREMER (GERMANY)
DR. M. NASTASI (USA)
PROF. H. NICKEL (GERMANY)
DR. NIELS HANSEN (DENMARK)
PROF H.R. OTT (SWITZERLAND)
DR. H. SCHUSTR (GERMANY)
DR. O.T. SORENSEN (DENMARK)
PROF. J.L. SOU9UET (FRANCE)
PROF S. BEDAIR (USA)
DR. N. EL MASRY (USA)
- v - SCIENTIFIC COMMITTEE (Continued)
II. NATIONAL
PROF. A.A. ABDEL AZIM (CMRDI CAIRO) PROF. A.R. ABDEL HALIM (CAIRO UNIV) DR. O. ABDEL WAHAB (INDUSTRY) PROF.A.AAMMAR (MENOF UNIV) PROF. S.R. ATALLA (CAIRO UNIV) PUOF. R. EL KOUSSY (CAIRO UNIV) PROF. N. EL MAHALLAWI (AIN SHAMS UNIV) PROF. M.A EL MOUSLY (AIN SHAMS UNrV) PROF. M EL OKKER (AZHAR UNIV) PROF. S.M. EL RAGHY (CAIRO UNIV) DR. A. EL SEBAI (INDUSTRY) PROF. A. EL SHANSOURY (AEA CAIRO) PROF. A. F. EL SHAZLY (AIN SHAMS UNIV) PROR. N. EL TAYEB (INDUST CONSULT) PROF. M. FARAG (AUC, CAIRO) PROF. M. FARID (CAIRO UNIV) PROF. N.M. GHONEIM (NRC CAIRO) PROF. M.M. HAFEZ (ASSUIT UNrV) DR. A. HELMI (INDUSTRY) PROF. R. KAMEL (CAIRO UNIV) PROF. S. KANDIL (ALEX. UNIV.) PROF. A. NOFAL (CMRDI CAIRO) PROF. A. RAMADAN (MINIA UNIV.) ENG. S. REDA (INDUSTRY) PROF. M. TAHA (AIN SHAMS UNIV) PROF. G.H. TALAAT (NRC CAIRO) PROF. B. ZAGHLOUL (CMRDI CAIRO) PROF. F.N. ZEIN (ALEX. UNIV.) VI - CO-SPONSORS
1. UNESCO, CAIRO, EGYPT 2. MINISTRY OF INTERNATIONAL COOPERATION, EGYPT 3. MINISTRY OF TOURISM, EGYPT 4. ICTP, TRIESTE, ITALY 5. ARAB ATOMIC ENERGY AGENCY, TUNISIA 6. ACADEMY OF SCIENCE & TECHNOLOGY, EGYPT 7. THE HOLDING COMPANY FOR METAL INDUSTRIES 8. MISR ALUMINIUM COMPANY, NAGAA HAMMADI 9. EGYPTIAN IRON AND STEEL COMPANY 10. EL NASR CASTING COMPANY 11. INVAP, EGYPT 12. EGYPTIAN ELECTRICITY AUTHORITY 13. ARAB ORGANIZATION FOR INDUSTRIALIZATION 14. ELECTROMITRY 15. TECHNICAL & CHEMICAL SERVICES CO. 16. JEOL EGYPT 17. SCIENTIFIC SERVICES CO.
- vii - STATEMENT OF HIS EXCELLENCY Eng. M.M. ABAZA MINISTER OF ELECTRICITY & ENERGY IN THE OPENING CEREMONY OF THE FIRST INT. SYMPOSIUM ON ADVANCES IN MATERIALS SCIENCE SAMS-94
DEAR GUESTS FROM ALL OVER THE WORLD
LADIES AND GENTLEMEN
IT GIVES ME A GREAT PLEASURE TO WELCOME YOU IN EGYPT ON THE OCCASION OF THE INAGURATION OF THE FIRST INTERNATIONAL SYMPOSIUM ON ADVANCES IN MATERIALS SCIENCE SAMS-94.
WE ARE HONOURED AND PREVILAGED TO HAVE WITH US THIS GATHERING OF MORE THAN FORTY DISTINGUISHED SCIENTISTS REPRESENTING UNIVERSITIES AND RESEARCH INSTITUTIONS OF 17 COUNTRIES FROM NORTH AND SOUTH AMERICA. EUROPE, ASIA AND AFRICA.
THE TOPICS ADDRESSED IN THIS SYMPOSIUM REPRESENT THE MOST RECENT ADVANCES IN SPECIAL MATERIALS COVERING CERAMICS. COMPOSITES, SUPERCONDUCTORS, SUPERALLOYS,
- viii - HIGH TEMP. MATERIALS AS WELL AS SYNTHESIZED AND MODIFIED SURFACES. THESE ARE JUST EXAMPLES IN THE VITAL AND RAPIDLY DEVELOPING .AREA OF ADVANCED MATERIALS.
THE INTERACTION WITH THE RECENT ADVANCES IN THE MATERIALS SCIENCE THROUGH HOLDING SYMPOSIA OF THIS TYPE, ATTRACTING SCIENTISTS AND EXPERTS OF THIS HIGH CALIBER SHALL HOPEFULLY HELP TO COPE WITH CHALLENGES TO DEVELOP MANY INDUSTRIES. THIS INCLUDES ENERGY INDUSTRIES : CLASSICAL, NUCLEAR AND RENEWABLE, IN ADDITION TO METALLURGICAL, CHEMICAL AND PETROCHEMICAL INDUSTRIES.
I AM AWARE ALSO THAT YOUR SYMPOSIUM IS DEVOTING SESSIONS TO DISCUSSING THE STRATEGY OF METAL INDUSTRIES IN EGYPT AND I AM SURE THAT THE PARTICIPATION OF THIS OUTSTANDING FOREIGN EXPERIENCE WILL BE VERY BENEFICIAL ONCE AGAIN I WELCOME YOU IN OUR COUNTRY AND WISH YOU A SUCCESSFUL MEETING AND A PLEASANT STAY IN EGYPT.
- IX - Firnt International Spring School & Symposium on Advances in Material Science (SAMS 94) Cairo, 15 • 20 March 1994
PREFACE
It gives me a great pleasure to introduce these valuable proceedings to the scientific community in the field of materials science in Egypt and in the world at large. These two volumes embody invited lectures and research papers from distinguished scientists in various disciplines of advanced materials representing reputed research institutions and universities all over the world. These scientists gathered in Cairo to contribute to this First International Spring School and Symposium on Advances in Materials Science (SAMS 94) in celebration of the 60th birthday of the AEA ex-chairman Prof. F.H. Hammad, a distinguished materials scientist. It gave the AEA a great honour to host this distinguished group of scientists, who contributed to the success of the symposium through their presentations and discussions, and who generously displayed their experience in vital advanced fields of materials science. This covered areas like composites, ceramics, powder metallurgy, superalloys, solid state chemistry, superconductivity.... etc. 1 would like to thank all contributors to the success of the symposium; the scientific committee, guest scientists from all over the world and local participants. I would also like to thank all members of the organizing committee, who worked extremely hard for almost one year to make this international event a real success. The effort given in the editing of these proceedings is greatly appreciated. y^Ut F fi\ Prof. H.F. Alz~-—± Chairman, AEA, Egypt January 1995 - x - First International Spring School & Symposium on Advances in Material Science (SAMS 94) Cairo, 15-20 March 1994 CONTENTS OF VOL. II Page 1- A Study of the Durability of Palm-Tree Fronds as a Reinforcing Fibre-Cement Composite Used for Low-Cost Construction Material. Samir Abdel-Azim M., Khedr S. andMorsyS. 9
2. Squeeze Casting of Hybrid SiIicon-12% Aluminium Matrix Composites. Khedre C, Attia AN and El-Desouky A.R. 25
3- Impact and Fracture Toughness Properties of Austempered Ductile Iron. Ahmed F.S. and Ghoneim MM. 45
4. Instrumented Impact Properties of Some Advanced Nuclear Reactor Pressure Vessel Steels. Ghoneim MM., Nasreldin AM., El-Sayed A.A., Pachur D. and Hammad F.H. 61
5. Fracture Analysis of Silicon Nitrides AbdEl-Razik ID., BishayAF., ElAslabiAM. KleistG.and Nickel H. 77
6. Formation of Austentic Stainless Steel by Solid State Reaction El-Eskandarany M.S. and Ahmed HA 93
7. Formation and Characterization of Nickel- Titanium Hydride Alloy Powder. El-Eskandarany M.S. and Ahmed HA 113 -1- Reactive Powder Metallurgy for NiAl-Base Alloy and Composites. Mohamed K.E., Stover D. and Buchkremer H.D 131
Mechanical Alloying of NiAl-Cr Eutectic Alloy with Non-Metallic Dispersoids Mohamed K.E., Stover D., Buchkremer H.P. and Hammad F.H. „..„ 157
The Aluminium-Oxygen Coordination of Anodically Formed Oxide Layers on Pure Aluminium. El-Mashri S.M. 173
Behavior of Ni and Fe Base Alloys in Arsenic Vapor Atmosphere. Afifi Y., Abd El-Razik L, Agamy S. and Krawczynski., S , 181
Crevice Corrosion of Hastelloy C-276 and Inconel-625 Alloys in Chloride Environments. I: Determination of the Critical Crevice Solution (CCS). GadM.MA., El-Sayed AA and El-Raghy S.M. 217
Crevice Corrosion of Hastelloy C-276 and Inconel-625 Alloys in Chloride Environments. II- Effect of Bulk Solution Environment Gad M.M.A., El-Sayed AA. and El-Raghy S.M. 239
A Study on the Corrosion Behaviour of X-750 Nickel Base Alloy. Abd-Elhady MA. and Kassem MA 255
-2- 15. Sulfidation of Iron-25% Managenese Alloy in H2S-H2 Environment at Temperature in the Range 973-1173 K. El-Refaie F.A., Ahmed H.A., Abdel Reheim N.A. and Smeltzer W.W. 265
16. Effect of Ultrasonic Agitation on Chromium Electroplating. Morsy S.M., Hosny A. Y., El-Rafey M. W. and El-Saycd MM. 285
17. Recrystallization of Cold Worked AI-AI2O3 Alloys and Commercial Purity Aluminium. Taha A.S., El-Houte S., Hammad F.H. and Hansen N. , 297
18. Effect of Type of Cooling on the Phase Transformation of Zircaloy-2 and Zirconium-lwt% Niobium. Taha A.S. 313
19. Microstructural Characteristics of Splat Cooled Al-Cu Alloys. Abbas A., Ramadan R.M., El-Nakhali A.E. and Ibrahim S. 329
20. Formation and Thermal Stability of Icosahedral Phase in Rapidly Solidified Al-15 Mn-2Ti Alloy. Ramadan R.M. , 339
21. Plastic Deformation Instability in Tension in Alloy 800 H. Abd El-Azim M.E. and Hammad F.H. 351
-3- 22. Evaluation of Abrasive Wear Resistance of Carburized and Plasma Sprayed 15 CrNi6 Steel. NasrH. and El-Ghazaly S.A 363
23. Characterization of Hetero-Epitaxial Structures by X-Rays and Electron Microscopy. El-Masry N.A 373
24. An Overview on Thin Films Prepared by Chemical Vapor Deposition. Abd El-Razik I.D., Afifi Y.K. and Krawczynski S. ... 389
25. Effect of Thermal Barrier Coating on Life* ime of FeCr Alloy Under Thermal Shock Conditions. Waheed A.F. and Soliman H.M. 417
26. Characterization of Indium-Doped Tin Oxide Films Prepared by Spray Pyrolysis. Afifi H.H.,Khalil M.S. and MomtazR.S 431
27. On the Electrical and the Optical Properties of CdS Films Thermally Deposited by Modified Source. AshourA., El-Kadry N. and Mahmoud S.A 441
28. Substrate Temperature Effects on Tin Oxide Films Prepared by Spray Pyrolysis. AfifiH.H., TerraF.S. and MomtazR.S 453
29. Effect of Temperature on Sn02 Thin Films Pyrolytically Deposited. MomtazR.S., DarwishS. and Afifi H.H. 463
-4- 30. Structural and Optical Propeties of CuO Films Prepared by Spray Pyrolysis. Afifi H.H., Nasser S.A. and El-Hakim S.A 483
31. Effect of Modified Evaporation Source on the Structural Characteristics of CdS Films. AshourA. and Mahmoud S.A 497
32. Advanced Dilatometric Technique "SID" for Studying the Kinetics of the Intermediate Stage of Sintering. Example: Sintering of U02 and Ce02. El-Sayed Ali M., Sorensen O.T. and El-Houte S 507
33. The Use of Alumina or Mullite in Fabricating Silicon Carbide Composite Materials. El-Masry M.A.A. and GadallaAM. 517
34. St able/Met as table Ce-TZP, Effect of Chemical Composition and Preparation Method. El-Houte S 535
35. Radiation Damage of Metals by High-Energy Charged Particles. Reutov V.F. 549
36. Swelling of Fe-30% Ni Alloy during Ar+ Ions Implantation. SoukiehM. andAl-MohamedAli 559
37. Evaluation of the Damage Introduced in Materials Proposed for the First Wall of Fusion Reactors Using TRIM code. Part I. Neutrons Effects. El-Shanshoury LA., Orabi G.I. and El-Shanshoury A.I. 569
-5- 38. Evaluation of the Damage Introduced in Materials Proposed for the First Wall of Fusion Reactors Using TRIM code. Part II. Alpha-Particles Effects. El-Shanshoury I.A., Orabi G.I. and El-Shanshoury A.J. 585
39. Study of Electrode Kinetics at the Interface Between High-Tc Superconductors and Solid Electrolytes. Zaghloul Hala S., Abdel RaoufM.A. andLorenz W.J. 599
40. Advanced Aerogel Materials for Photocatalytic Detoxification of Cyanide Wastes in Water. Ahmed M.S 609
41. Measurement of the Thermal Properties of Co(N03)2 and Acetamide by a Modified AC-Heated Wire Technique. Atalla S.R., Attia G.A., El-Naggar M.M., El-Sayed T.A., El-Sharkawy A.A., N.K. Mina and Kenawy M. 623
42. Measurement of the Thermal Properties of AgNOg and Cd(N03)2 by a Modified AC-Heated Wire Technique. Atalla S.R., El-Naggar M.M., El-Sayed T.A., El-Sharkawy A.A., Mina N.K. Gamal G.A. andNagatAT. 633
43. Evaluation of the Efficiency of Hydrophobic Agents by Using Neutron Radiography. Scherpke G. and Ashoub N. 643
• 6- 44. Neutron Diffraction and Mdssbauer Measurements on Gallium Substituted Zinc Ferrite. Fayek M.K., Sayed Ahmed F.M., Abu-Hclal II., Mostafa M.F. and Kaiser M. 651 45. The Determination of Cu/Ni Ratio in Copper-Nickel Alloy by Neutron Capture Gamma-Ray Method. Abulfaraj W.H., Abdul-Majid S. and TuruustaiiL 0 667
46. Mossbauer Invistigation of A Canted Spin Fe(II) System. Moustafa M.F., Abd El-Kader MM., Atallah S., Soliman S., El-Esawi M., andEmrick R 675
47. Infrared Absorption Spectra of Egyptian Serpentine Rock. KhalU E.M.A 685
48. Infrared Absorption Spectra of Silica By-Product of the Egyptian Ferro-Silicon Factory. Khalil E.M.A 692
49. Relation Between Maximum Degree of Hysteresis and Both Temperature and Initial Pressure During Hydriding of LaNi5 Isaack S.L., Karakish E.A. and Shaaban H.I. 701
50. Phase Transitions During Hydriding Activation of LaNig Intermetallic Compound. Isaack S.L., Karakish E.A and Shaaban H.I. 711
Author Index 721
-7-
First International Spring School & Symposium on Advances in Materials Science (SAMS 94) 15-20 March 1994
A STUDY OF THE DURABILITY OF PALM-TREE FRONDS AS A REINFORCING FIBRE - CEMENT COMPOSITE USED FOR LOW COST CONSTRUCTION MATERIAL
M. Samir Abdel-Azim', S. Khedr*\ and S. Morsy"' • Quality Assurance Dept., Nuclear Safety Center, Egyptian Atomic Energy Authority ** Construction Engineering Dept., American University in Cairo *** Metallurgy Dept. Nuclear Research Centre, Egyptian Atomic Energy Authority
ABSTRACT PTF, as natural fibres, have many good characteristics that make it suitable for the enhancement of the characteristics of cement-based composites to be used economically in the construction of houses. In the present work, physical, mechanical and chemical tests were carried out to examine the durability of PTF. Also the effect of adding PTF leaf fibres to concrete was stuiied and was found to be suitable to improve the tensile strength and workability of concrete.
I. INTRODUCTION The addition of a random dispersion of small fibres to brittle materials as cements and concretes significantly improves the strength, ductility, resistance to cracking and durability of the resulting composite materials (1), (2), (3). The effect of adding the fibres is to increase the local tensile Strength of the composite material almost at all points in different directions, hence preventing the spread of cracks. The fibres have also the effect of evenly distributing drying shrinkage. So that in this case and in the case of direct or flexural tensile stress, a number of very fine cracks distributed all over the member will develop instead of having one or two major cracks with a very appreciable crack width (4). The absence of the transverse cracks in the tensile zone of the fibre reinforced cements will increase the rupture strength ot such material. In other words, the addition of fibres to concrete, is to convert the sudden brittle failure of plain concrete in tension, into a gradual and a lot more ductile failure, an advantage which would be appreciated in design of building houses or schools in regions with earthquakes activity. -9- During the past 30 years, and recently (1) (2), substantial interest has been shown in fibre reinforcement of concrete, by research organizations and construction industry, after the pioneering research of Ramualdi and Baston (5). on steel fibre reinforced concrete was published in 1963. Since then considerable development have taken place, especially with steel fibres. However the rusting of reinforcement in concrete is the single most important cause of poor durability of structural concrete (1) (6) (7). Since concrete cracking occurs because ihe corrosion product of steel, hydrated red rust, has a volume more than four times that of the metallic iron from which i' was formed. The forces generated by this expansive process can far exceed the tensile strength of the concrete with resulting cracking (6) (7). Steel corrosion not only causes cracking and revelling of the concrete but may also cause structural failure resulting from the reduced cross-section and hence reduced tensile capacity of the steel. This form of damage can be a significant problem in aggressive environments such as petroleum industry regions (6). The traditional approach adopted by design codes for overcoming the problem is to use large covers, with relatively high cement contents, to limit crack widths. However, increasing the cover while still maintaining a limiting crack width is self-defeating, resulting only in an increasingly limited stress in the reinforcing steel and a lowered load- carrying capacity (1).
From durability point of view, steel fibres are susceptible to the problem of corrosion at cracks, particularly in areas where sulphur pollution is already an existing problem (8) (9). Because of this, new fibres and composites are continually being developed. In particular, the fibres being used other than steel, are carbon fibres, glass fibres, plastic fibres and organic fibres. Cook (11) shows that natural fibres of vegetable origin such as straw, grass and reeds have been used to reinforce mud bricks as early as 1400 B. C. as found for example near Baghdad (Ziggurat) built in 1400 B. C. Even to-day, straw is still being used in some developing countries to reinforce mud plaster and bricks.
Another example of composite materials which have been used almost all over the world for more than 75 years is asbestos-cement. However, it is now evident that asbestos has a hazardous effect on health, and efforts have been made to find a convenient substitution to replace it. In fact, there has been a resurgence of interest in recent years in vegetable fibres in developing countries and elsewhere. This had been brought about by the energy crisis, a general oversuppiy and hence relatively cheap source of fibres and the need to reduce foreign reserve expenditure on importing other fibres. In Egypt and many other countries, PTF are very available, as natural fibres. -10- Our preliminary investigations which were published and presented in The American Concrete Institute Magazine (12) and in The International Seminar, "Building The Future" (13), and the great interest from the Construction Research Establishments to replace steel reinforcement with cheap and efficient fibres in building economic houses, encourages us to continue our studies of the durability of PTF to use it efficiently as a low-cost building material.
The main purpose of the present investigation is to determine if PTF exposure to chemical solutions of different compositions with pH • 1 to pH= 13 and cement solutions, would lead to chemical deterioration or would affect PTF strength if used as a construction material,
II - EXPERIMENTS AND PROCEDURES II. 1. Resistance of PTF to Chemical Deterioration: To determine the resistance of PTF to acidity and basicity, separate sections of dry PTF were exposed to various chemicals and cement solutions for different periods and a variety of temperatures. Initially the leaves were removed by clipping as close as possible to the stem without damaging the stem. The stem was then cut into sections of approximately 6-7 cm in length. Again care was taken to minimize any damage to the outer surface of the stem. Each section was wiped with a tissue to remove any dust accumulated 011 the surface. The clean sections were immersed in 100 ml of solutions with pH - 1 to pH = 13 at different temperature up to 64'C. The sections were allowed to remain in contact with the solution for 1,6,12,18 and 24 hours. The pH was measured after the required treatment. The samples were weighed before and after each treatment and after each section was allowed to dry at room temperature for 48 hours and after 28 days. Table 1 includes overall results of ail PTF soaked in different chemicals.
11.2. To find the effect of cement on PTF, chemical deterioration and on its strength, mixtures of 50 grams of cement and 100 ml of deionized water were prepared. Thirty cm rods of PTF were fully immersed in the mixture in tubes that were sealed and shaken mechanically for 1, 6, 12, 18 and 24 hours in different solutions at 28*C, 52' and 64*C. Measurements of pH were obtained before PTF immersion and at the end of the experiment. Samples were weighed at the beginning of the experiment, at the end of each of the periods mentioned above, after 48 hours and after one month. Since the purpose of this work is to determine if PTF exposure to chemical solutions of different composition and under certain conditions would lead to chemical reactions that would cause weight loss (chemical deterioration) or would affect PTF strength if used as a construction material, therefore data related to weight were treated as follows Weight was measured at -11 Table 1 A = Height before experiment, B = Wt, 48 hours after the end of experiment. C = Wt. after 7 days, D = wt. after 28 days, R = Reagent, S = Soaking duration in hours
1 A d C 0 8 1 Experiment IT- 26.3»C 5.121 5.14 5. 19:-" 5.15 1 1 R - Sulfuric Acid 4.53 4.68 4.71 4.58 12 pH- 1.08 3.18 3.21 3.27 3.20 24
T - 28. VC 7.14 7.06 7.23 7.20 1 R • Hydrochloric acid 11.51 11.40 11.93 11.71 12 pH- 2.68 5.73 6.61 6.97 6.81 24 T » 25.7»C 4.56 4.58 4.91 4.78 1 R - Sulfuric Acid 6.21 6.28 6.66 6.52 12 pH- 3.06 4.48 4.57 4.50 4.41 24 T - 24.7«C 7.08 7.22 4.17 7.08 1 R - Sodium phosphate 4.12 4.35 4.21 4.12 12 pH- 5.46 3.95 4.21 4.15 4.09 24
T - 24.8»C 5.83 5.80 5.93 5.82 1 R » Sodium phosphate 6.72 6.66 6.89 6.70 12 pH- 7.64 4.12 4.01 4.30 4.20 24
T - 25.2»C 3.18 3.17 3.28 3.18 1 R * Sodium Hydroxide 6.29 6.29 6.33 6.23 12 pH- 9.41 6.08 6.18 6.06 5.97 24 T - 27.7'C 9.64- 9.68. 9.78 9.69 1 R - Calcium Hydroxide 10.11 10.51 10.42 10.20 12 pH- 12.51 7.33 7.53 7.50 7.27 24
-12- the end of the experiment, after 48 hours and after one month. The percentage of weight change was calculated for all periods in all experiments,
weight after exp. - weight at the beginning of the exp. %a = xlOO weight at the beginning of exp.
wt. after 48 hrs. of exp. - wt. at the beginning of the exp. %b = x 100 weight at the beginning of exp.
wt. after one month of exp.- wt. at the beginning of the exp. %c« : xlOO weight at the beginning of exp. Table 2, includes the results of PTF soaked in cement solutions at different temperatures (I. at 28*C, II at 52*C and III, at 64*C). 11.3. The Effect of Cement Solution on the tensile strength of PTF, To investigate the cement effect on the tensile strength of PTF, two groups (A and B) of PTF samples were taken from the same stalk and the same location on the stalk. Since it was found in our previous investigations (12,13) that the ultimate tensile strength of the peripheral portion of the stalk is higher than that of the core (this is because the peripheral fibres in the stalk are more dense and congregated), therefore, care should be taken to have specimens cut from the same stalk and the same location on the stalk. The total number of samples obtained was 14. The samples of group B were tested to determine the ultimate tensile strength of the as-received specimens. The samples of group A. which were treated with a cement solution for various lime durations and at room temperature, were also tested to determine its ultimate tensile strength. The results are shown in table 3. It has been noticed that there was an increase in strength in all the samples which were treated in cement solution. The strength increase percentages in core specimens were much higher than peripheral ones.
11.4. The Investigation of the Effect of Randomly Adding PTF Leaf Fibres to Concrete The variation of the fresh and hardened concrete behavior with the amount and length of added PTF leaf fibres was investigated. Tests conducted, in this stage on fresh concrete was the "Slump test", to measure its workability. For hardened concrete, the "cylinder or cube compressive strength" and the "split cylinder" tests were carried out -13- Table 2 X- Results of PTF Treated with Cement Solutions at 28*C ( Roost Temperature )
Exposure weight Change Percentage Tine (hours) *a Vb %c | 1 19.13 1.80 2.79 6 24.02 3.«1 3.05 12 52.91 9.43 2.79 18 78.19 14.89 2.29 24 104.98 16.93 2.18
ZJC- Results of PTF treated-with Cement Solution at 52*C
Exposure Weight Change Percentage Time (hours) %a %b %c 1 14.98 1.36 0.35 6 I 114.94 3.97 0.30 111.17 2.61 0.81 12 112.88 8.79 0.84 18 94.68 11.47 0.31 1 24
-14- Table 2 (Continue) III- Results of FTF treated with Cement Solutiona at 64*C
Exposure Weight change Percentage Time (hours) %a %b %c 1 29.02 2.60 0.43 6 100.96 16.93 0.64 I
12 144.40 29.91 0.09 18 126.93 9.22 0.12 130.23 40.32 0.83 "
veight after exp. - veighl at the beginning ol the exp. « n- — 2 !00 •weight at the beginning ot exp
vt. niter 4X hrs. nf exp - wt at the beginning or the e*p \b X 100 weight at the beginning of ezp.
•vi after one month of exp.- wt at the beginning of the cxp. r. c- 100 weight at the beginning of *xp.
-15- Tabla 3 UXtiaata Tonsil* Strength values for samples: aroup A: Xftnarsed la ceaent Solutions Group B i As - raeaived
Saaple 0T3 (H/BU2) Strangta Increase % Location No. aroup A Group B 1 105.7 88.3 19.7 Cora 2 116.5 92.3 26.2 119.3 103.1 15.7 100.8 81.4 21.8 238.5 224.5 6.2 Paripharal 220.1 207.8 5.9 207.5 191.1 8.5 240.8 221.7 8.6
Table 4 (a, b) a) Concrete Hl> Sroportione
Constituents .. % by vaighC Coarsa Aggregates 44 Fine aggregates 28 Cenant 18.8
Ha tar 9.2
b) Sradation of Coarsa Aggregate
Sieve size % passing h- 100 3/8" 60 No. 4 40
No. 8 20
No. 16 5
-16- to determine the compressive and tensile strength The relative quantities of the materials used in the concrete mix and the gradation of coarse aggregates in it are given in table 4 (a, b) respectively A cement aggregate ratio of 1 4 was used while a water / cement ratio of 0 48 was chosen and maximum aggregate particle size was 1/2 " PTF leaf fibres of approximate width of 15 - 2 00 mm and the chosen length was mixed with concrete. Three different values for the approximate fibre length were tried 12 cm. 6 cm and 3 cm. The fibre weight ratio was chosen to be 1%, 0.5% and 0.25% of the total weight of concrete.
ft must be noted here, that during a preliminary split cylinder test, it was noticed that in most ca.se^, concrete around the fibres was not completely hardened which may have caused a reduction in the resulting fibre-concrete strength. Hence, fibres were submerged in a solution of 50 gm of cement / 100 ml water for one hour before mixing it with concrete. This is to reduce water absorption by fibres during and after mixing with concrete which could affect the net water / cement ratio and at the same time, to strengthen the fibres itself by impregnation with cement solution, as already resulted earlier.
Concrete cylinders were cast, compacted using table vibrator and left in molds for two days. The specimens were then taken out and cured for 28 days in water tanks. II.5- Tests Performed on Fresh and Hardened Concrete Specimens. li) Slump test: It is the universally known test to measure the consistency of concrete and it was performed in accordance with ASTM-C 143,Slump of Portland Cement Concrete. (ii) Compression Test: Concentric compression force was applied to the ends of the concrete cylinder, diameter 100 mm, at constant rate of 9KN/sec, until failure happened at a load Lf, the compressive strength 4Lf 2 fc = N/mm 314 x (100 )2 (iii) Split cylinder test: This test is used to indirectly measure the tensile strength of concrete. The test was chosen here because of the suitability of the test arrangements and the difficulty of using the direct tension test on fibre - concrete specimens. A cylinder with diameter D = 100 mm and a length 1 =200 mm.is placed with its axis horizontal between the plates of the testing machine as shown in Figure 2. An applied load was increased at a rate of 9 KN/sec. until failure occurs at a load Lf. The failure takes place by splitting of the cylinder along the vertical plane connecting the two contact -17- /A i ir ~P
Split Cylinder Ten figure 1
Photo (a) Tensile - Strength Splitting test
Photo (b'> Compressive Strength test Figure 2 Pboio: J. bi of Ductile Failure of PTF Fibre-.-•;:3i'o,-:»(i concre1.- soec:mens. -18- lines vith the machine heads The tensile strength, based on linear elastic assumption. is calculated as:
4Lf
l'Sp ' N/mm* 3 14 x(100)(200) Table 4, shovs results of the slump test, compressive strength test and the split cylinder test of PTF fibre - concrete and a plain concrete Table 5 shovs results of the slump test, compressive strength test and the split cylinder test of PTF fibre - concrete in vhich the fibres •were immersed in cement solution fnr one hour before mixing. II.6. Palm Tree Frond Composition The purpose of this section istb study the reason behind the increasing ultimate.tensiie strength .of PTF after immersing in cement solution and to investigate its internal structure and composition. The PTF consists of a2 m to 5 m stalk that carried ap-tQ 800 mm Jong and 35 mm vide leaflets. The stalks are long, unbranched and filled inside vith cellular tissues vhich are traversed by strands of fibro-vascular bundles, vhich provide fluid transport apd structural stability,
Several thin sections oil the.PTE stalk vere examined using a phase contrast compound microscope vith mlcrophotographic facility. Good anatomical details under various illumination and magnification vere obtained. Figure 4 shovs the anatomic structure of cross section of PTF stalks, schematic but not to scale.
III. RESULTS AND DISCUSS IONS III.I. PTF Composition • The main features of PTF stalk, as shown in Figure 4, include an impermeable cuticle vith chloroplasts. to protect the stalk from dehydration, followed by a single layer of epidermis, large number of paranchyma cells and scattered bundles of fibro - vascular tissues of different sizes. The fibres combinedly form fibro - vascular bundles vhich provide fluid transport and structural stability. The peripheral bundles are small, compact and congregated The inner bundles are large and scattered. Nov, ve can rationalize the high tensile strength obtained for the immersed stalks in cement solution. As the cement solution can go through the fibro - vascular bundles and cementize it becomes more congregated, and this strengthens the specimen Moreover, the data presented in Table 3 demonstrate that the strength increase percentage vas much higher for core specimens than peripheral specimens. That is because the peripheral bundles are -19- Figure 3. Schematic representation of a Palm Tree Frond (PTF)
FIGURE 4. Anatomic Presentation of Cross-Section of PTF Stalk Schematic but not to scale, a) Cuticle b) Epidermis c) Fibro-Vascular bundles d) Parenchyma Cells -20- Table 5 Results Obtained Using as received fibres (without innersing ia Cement Solution*)
(JIIP Hi 1 III1M1HI Fibres Slump Average Average fc fsp 1 Mix No. Wt. % Average mm N/mm2 N/DB2 Length, mm 1 Plain Concrete 40 25.78 2.55
2 1 120 33 6.12 2.31
3 1 60 27 6.83 1.48
4 1 30 18 14.83 2.23
5 0.5 120 28 15.32 2.51
6 0.5 60 36 18.04 2.85
7 0.5 30 27 17.37 2.71
8 0.25 120 38 17.45 3.01
9 0.25 60 40 21.87 3.02
10 0.2S 30 39 21.02 2.42 .. Tabla C fteaulta Obtained Oaiag PTF Fibres Xmaersad in caneat Solution for oaa Bour Bafora Mixing with eonerata
Fibre* slump Average {Average fsp XiZ NO. wt % Average tr/nstt N/mB2 Length, ma 1 Plain Concrete 40 25.78 2.55
2 1 120 35 15.18 2.16
3 l 60 25 15.37 2.23 4 l 30 18 18.77 2.28
5 0.5 120 40 19.83 2.95
6 O.S 60 43 20.08 2.87 7 0.5 30 37 22.20 2.81 8 0.25 120 43 21.90 3.18
9 0.25 60 55 23.68 3.19
10 0.25 30 50 24.70 2.62 -21- small, compact and congregated, but the core specimens have more opened paths of wider vascular bundles, as shown in Figure 4, that enables the cement solution to flow through it.
III.2. The Effect of Randomly Adding PTF Leaf Fibres To Concrete: Results of fresh concrete slump test are summarized in Tables, 5 and 6. It is noticed that generally workability, with the slump as a measure, is improved as the amount of fibres is reduced. It may be concluded that for fibre - concrete composite vith good workability and uniform fibre distribution, the recommended fibre weight percentage may be Jess than or equal to 0.5% with a fibre average length in the range of 30-120 mm. Concerning the hardened concrete, it is reported in literature (4) that the addition of a random dispersion of small fibres to concrete and cement could be advantageous in improving the tensile strength, ductility, resistance to cracking and durability of the resulting composite materials, moreover, the present work adds new information, namely that more tensile strength of concrete could b« obtained by immersing the fibres in cement solution for one hour before mixing vith concrete as shown in table 6. Basically, the effect of adding the strengthening fibres is to increase the local tensile strength of the composite material at a large number of points, in different directions, depending on the amount of added fibres and on the random directions of these fibres, as clearly shown in the photo of the splilted cylinder and the photo of the compressed cube of figure 3- Hence, the existence of the fibres will "trap" cracks and stop or at least delay their spread, especially if those fibres were immersed in cement solution before mixing with concrete, to have better bonding with the concrete matrix. Concerning the compression test results which are summarized in tables 5 and 6, it is clear that the addition of fibres has no positive effect on the compressive strength of concrete. However, it was noticed that a relatively high compressive strength was obtained when the fibres vere immersed in cement solution for one hour, before mixing with concrete. The drop in the compressive strength of plain concrete by adding fibres to it may be attributed to the following: i- As the fibre percentage and length are increased, the probability of these fibres to nest together leaving large voids in the concrete is greater. This reasoning explains the very low value obtained for mix No. 2 with fibre percentage of 1% and fibre length of 120 mm. in table ii- It has been noticed that in many cases concrete around the fibres was not completely hardened In some cases, the ingredients, particularly cement, are found in a powder form. One explanation of this phenomenon may be that fibres alter mixing with concrete have absorbed more water denying cement around it enough water l'or the hydration process. The solution to this problem that is used in this investigation is: the fibres should be immersed in cement solution for one hour before mixing with concrete. The results of this solution can be noticed clearly by comparing the average compressive strength of mix No. 2 in Table 5 and table 6, which has been increased by more than twofold. iii- Fibres in large quantities seem to fundamentally alter the nature of cementitious matrices. When, a large amount of small diameter, uniformly distributed fibres are employed, it appears that the inherent tensile strength and strain capacity of the matrix itself is fundamentally enhanced. The presence of fibres can possibly arrest microcracking in matrix and delay localized micro cracking. Conclusion As conclusion, using PTF as fibre composite offers unique advantages for solving civil engineering problems in areas where conventional material fail to provide satisfactory service life. Unlike steel, a very cheap material of PTF are unaffected by electrochemical deterioration and can resist the corrosive effects of acids, alkalies and similar aggressive materials of different pH and under a wide range of temperatures as table 1 and table 2 demonstrate. The incorporation of PTF in cement based matrices resulted in producing inexpensive, lightweight, heat insulating building material that can sustain excessive strains without exhibiting sudden failure. An advantage which would be appreciated in design of building houses or schools in regions with earthquake activity. REFERENCES (1) Clarke, J. L„ "Fibre Composites For The Reinforcement of Concrete", Proceedings of "Building The Future" Seminar, Building Research Est., held at Brighton, U.K., April (1993). (2) Takayuki Hiral, "Use of Continuous Fibres for Reinforcing Concrete, High Performance Durability". Concrete International, American Concrete Inst. Magazine. Vol. MNo. 12 December (1992). (3) ' Shah, S. P., "Fibre Reinforced Concrete Application", Concrete Yesterday, Today and Tomorrow, Concrete International, Am. Concrete Inst. Magazine, Vol. 12, No. 3. March (1990). (-4) Krenchel, H. and Jensen, H. W„ "Organic Reinforcing Fibres for Cement and Concrete", Proceedings of the Symposium on Fibrous Cements, Held in London on 16 April, (1980).
-23- (5) Ramauldi. J. P. and Baston. G. B., "Mechanics of Crack Arrest in Concrete", Proceedings of the American Society of Civil Engineers. Vol 89,No.EM3,p.pH7,June(l%3). (6) Cady. P. D., "An Overview of Concrete Durability Programs In The Arabian Gulf Region", The Arabian Journal For Science and Engineering Vol. 11. No. 2. p p. 120 (1985) (7) George J. Verbeck, "Mechanisms of Corrosion of steel in Concrete. Vol. 24, No. 11 p. 21 (1986). (8) Gany Rao U.V., and Faza, S, S., "pre - and Post - Cracking Deflection Behavior of Concrete Beams Reinforced with Fibre". Proceedings of "Building The Future" Seminar, Brighton, U. K., April (1993). (9) Grzybowski, M and 5hah, S. P.. "Shrinkage Cracking of Fibre Reinforced Concrete", ACI Materials Journal, Vol. 87, No. 2, March-April (1990). (10) Sakai H. et al., "Experimental Study on Flexural Characteristics of Carbon Fibre Reinforced Cement Composite".Proceedings of "Building The Future". Seminar held at Brighton.U. X., April (1993). (11) Cook, D. J., "Concrete and Cement Cmposites Reinforced with Natural Fibres", Proceedings of the Symposium on Fibrous Cements held in London on 16 April, (1980). (12) Abdel-Azim. M. S.. "Palm Tree Fronds For Concrete Roof Reinforcement", Concrete International, The Am. Concrete Inst. Magazine, Vol. 14, No. 12 December (1992). (13) Abdel-Azim, M. S„ "Application of Palm Tree Fronds in Building Heat-Insulating Houses", Proceedings of "Building The Future", International Seminar held at Brighton, U. K„ 19 April (1993).
-24- First International Spring School & Symposium in Advances in Material Science (SAMS 94) Cairo, 15-20 March 1994 SQUEEZE CASTING OF HYBRID SILICON-12% ALUMINIUM MATRIX COMPOSITES O.M. Kherdre, AN. Attia and AR. El-Desouky Prod. Eng. and Meek. Design Dept., Faculty of Eng., Menoufia Univ., Egypt.
Abstract The squeeze casting technique is becoming more and more established as the major liquid metal process for fabrication of metal matrix composites (MOMS). It combines near-net shaping of parts with a high process speed and controlled microstructure engineering. The present work is concerned with the study of the effect of squeeze casting on the properties of Al-12% Si alloy reinforced with S^C particles in addition to graphite powder. The latter was added to improve the micro structural homogenity as well as the uniform distribution of SjC with the resultant decrease in the preform compression. The added graphite and S^C -contents were up to 5 and 15 wt. % respectively. The composite specimens were squeeze cast under 50, 80 and 125 MPa pressure applied for a duration of 120 S. The experiments were carried out under conditions as close as possible to these of traditional pressure die casting. To characterize the mechanical properties of the composites the tensile strength has been measured because it is an important material parameter, and relevant to the understanding of bonding between particles and matrix. In addition, electrical resistivity measurements were made to follow the age hardening capability of composite materials. SEM observations were performed to
-25- characteriste the frature surface as affected by the aforementioned variables. Introduction Aluminium and its alloys provide designers and engineers with a group of materials with valuable properties: high strength/wt-ratio; excellent corrosion reistance; high electrical and thermal properties; and ease of fabrication. Cast Al-alloys have successfully replaced ferrous casting in transport-engineering application enabling fuel economy to be gained through weight reduction. In many applications requiring resistance to wear, Al-alloys have proved effective. the hypereutectic alloys, containing 14-20% Si, have been used for automobile blocks monocoque construction. The Reynolds Metal Company (USA) developed an alloy containing 16-18%Si, designated . 390 alloy, which successfully used in the 1970 Chevrolet Vega 2300 engine. The use of an etch treatment to leave the Si-particles standing proud of the load - bearing surface was the first success, making it possible to omit a cast iron liner. In Europe, cylinder blocks without cast iron liners have been increasingly used, for ex. the Reynolds 390 alloy is used in the Porsch 928, the Diameter - Benz 450 SLC, and the Rolls - Royce Camargue. Considerable economic savings are possible if the Al alloy surface itself can be used directly as the bearing surface. A more recent development has been in the production of Al-graphite alloys, which show very interesting tribological behaviour1 . These alloys show every sign of further significant development, the presence of graphite particles in an Al alloy matrix improves its resistance to seizing, resistance to scoring wear, machinability and damping capacity'- . Hypereutectic Al - Si alloys containing dispersed graphite particles are suitable for cylinder or cylinder sleeves in internal combustion engines^.
.26- Preparation of Al-graphite particle composites by conventional melting and casting technique is difficult, since graphite is not wetted by liquid Al below 1353 K'4'. This leads to rejection of graphite particles from the melt. Two main approaches have been employed in the last few years to improve the wetting between graphite and liquid Al. One of the method in the use of Ni - or Cu - coated graphite particles . But copper or nickel coating of graphite particles is expensive and may adversely influence the properties of the resultant alloy. The other is mixing of graphite particles with liquid Al using ultrasonic vibrations' . However, the use of an ultrasonic probe in the large - scale production is prohibitively expensive. Infiltration of coated graphite particles by molten Al alloy was kept patented' . However, the methods mentioned above require special treatments for either the melts or graphite particles. Consequently,, the processes are complex and time consuming and therefore unlikely to be suitable for mass production. The driving force for the development of MMC has been the quest for faster, more fuel efficient airplanes. Indeed, in a typical commercial aircraft for every kilogram of weight reduction $550 savings in fuel costs are projected for the life of the aircraft at todays fuel prices'8'. While fiber reinforced composites have some very exciting properties, their prices make them prohibitive for all defense or very specialised applications. For example the cost of boron fibers is $/kg.660 When this is compared with the price of particulate SiC of $/kg, 4.85 it becomes clear that PMMCs have more appeal for general applications, in particular for automotive applications' . It should be further noted that the fabrication price for PMMCs is significantly lower than for fiber reinforced MMCs. Al-based MMCs with SiC reinforcement are currently being produced by powder, spray and cast routes'10'. These materials are being evaluated for a range of applications primarily in the aerospace sector' \ By combining adequate
-27- matrix and reinfrocement materials, it is possible to produce PMMCs with special properties. One example is the development of a three - component composite consisting of Al-matrix and two different reinforcement, SiCp and elemental Si'"'. This composite has a very low coefficient of thermal expansion while preserving good thermal conductivity and is thus an excellent material for electronic packaging. The present work is concerned with the study of squeze casting of Al-12% Si alloy matrix containing graphite and SiCp dispersions. The SiCp - addition has the aim to improve the strength properties and thus to compensate for the drop in strength due to graphite addition. To assess the properties of the obtained hybrid composites serveral tests were performed in addition to fractographic analysis. Experimental Procedure Fig {1) presents the flow chart of the experiment. All the experimental procedures described in this section are referred to this chart. Al-Si eutectic alloy of a nominal composition 12 wt% Si was first prepared by mixing the appropriate amounts of 99.7% commercial purity Al with a 50% Al-Si master alloy. The hybrid composites containing silicon carbide, and graphite particles were prepared using the set-up shown in Fig. (2). A single size range of uncoated synthetic graphite particles was selected based on previous experience, i.e. 75-250 um diameter of particles. Particles smaller than 75 mm led to undue segregation and agglomeration during casting, wheras particles larger than 250 um were believed to have no favourable effect on mechanical properties. The weight percents of graphite used in the present hybrid composites were 3 and 5%. This was because the composites having graphite less than 3% would not posses adequate wear resistance' , on the other hand more than 5% graphite
-28- f-0
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.29- would lead to rejection from the melt. A single particulate size range, 75-180 Jim, of SiCp was selected to yield 5,10 and 15 wt% of the hybrid composites. To increase the wettability of graphite and SiCp particles by Al-Si alloy, magnesium (0.5%) was added to the melt along with graphite and SiCp particles. It was reported that Mg has a good influence on the bonding between matrix and reinforcement^2.'. SiCp and graphite particles were preheated at 800°£ and added to the melt through a vortex created by an electrical stirrer (500:600 r.p.m.). Mixing was continued to obtain a uniform distribution of SiCp and graphite. The pouring temperature of the melt* and the die temperature were maintained at 710*C and 500*C respectively. The viscous slurry was* ?cBntt4) 30- K-HHh i were machined from the squeezed composites with the guagc length of specimens parallel to the longitudinal axis of the casting. Testmetric Machine WP -sed throughout the present work for room-and high-L,mperature. Tensile testing was performed at constant cross-head speed of !jmm/min. A minimum of three specimens were tested for each case. Hardness Tesli> The Rockwell hardness B of the composites was established using a hardness tester (Akashi) model (AR-10). A minimum of three hardness readings on the diameter and six readings along the length were taken for each specimens. MctaU.ogra.phic Examination and SEM Fractograpky Specimens of cross section of 25x15mm were cut from different parts of the casting and polished for metallographic examination. Specimens were viewed using Metaval - metallograph and representative photographs were taken. For the fracture surface examination, the specimens were carefully machined from the broken tensile specimens. The fracture surface were cleaned ultrasonically and provided with silver coating. The fracture surfaces were examined using JEOL 840 A scanning electron microscope. Resistivity Measurements The resistance of the samples was determined using the conventional four - prnhe method using the circuit shown in Fig. (3). In these measurements silver paste was used as electrodes. DC programmable current source "Keithley, Model 244" was used with applied current of 10 mA. The possibility of the disturbing thermopower was eliminated by changing the current automatically. The voltage drop was detected using an accurate digital nanovoltmeter "Keithley Model 181". All the system was computerized - When an electric tli passed through the sample, resistance R was -31- Fig. 4 Microstructure of Al-12% Si alloy a|Casted in metal mould, 400x. b)after T6 heat treatment, 300x c) squeeze formed at 80 MPa, 400x a) b) Fig 5 Microstructure of: a) Al-12 Si/5 SiCp, 80 MPa, 400x b) Al-12 Si/5 SiCp- 3Gr, 80 MPa, 400x mmm^m Fiug 6 Photogr^^h showing macrostructure of squeeze formed ingot of hybrid composite (Al- 12Si/10 SiCp- 5 Gr), 80 MPa -32. obtained using Ohm's law. The resistivity (p) was calculated from the following equation; p j. 2 7i a Q^'c where b is the specimens thickness (cm). Results And Discussion: Microstructure In the Al-Si alloy system; the eutectic forms at 12.6 wt% gj[l3] r^he structure of the investigated matrix alloy (containing 12 wt% Si) solidified in a metallic mould is shown in Fig. (4-a). The structure consists mainly of primary Al dendrites with an average dendritic arm spacing on the order of 25 um and eutectic silicon in the interdendritic region and around the dendrites. The eutectic Si is plate like in appearance and some of these plates are interconnected. The effects of T6 heat treatment (solution heated at 520°C for 8th, quenched in warm water at 60°C and aged at 180°C for 6h) and squeeze casting on the structure of the investigated matrix material are shown in the photomicrographs of Fig. 4-b & -c respectively. The heat treatment of Al-12 Si alloy resulted in a significant change in the morphology of eutectic Si from plate like to nearly spherical. This change in morphology of Si due to heat treatment also results in an increase in Rockwell hardness as will be mentioned later. Stirring of the melt followed by squeeze casting resulted in a breakdown of the dendrites with a considerable decrease in the dendritic arm spacing to approximately its half value. In addition, stirring and squeezing changed the morphology of the eutectic Si from plate like to the globular Form Fig. (4-c). The microstructure of Al-Si/SiCp and Al-Si/SiCp + graphite (hybrid) composites are shown in Fig. (5). The introduction of graphite particles, Fig. (5-b), offers the possibility of producing composites at a relatively lower pressure with a good property profile. Homogenity is an important factor in the production of Al/Sic composites. It -33- can be increased by decreasing SiCp-content, by changing the squeeze casting procedure and by introducing graphite particles, the uniformity of distribution is demonstrated by the macrostructure of the ingot casting shown in Fig (6). Furthermore, excessive chemical reaction between SiCp and the reactive Al melt can be avoided through the addition of graphite and application of pressure during solidification. This effect is of special interest, since A14C3 is a brittle intermetallic compound and results in the degradation of the SiC reinforcement. Moreover, A14C3 is soluble in water and its formation during the fabrication process increases significantly the viscosity of the melt which can make casting of the composite impossible if produced in sufficient quantities. This is why efforts have therefore been made by many groups to monitor as quantitatively as possible the extent of AI4C, formation. Mechanical Properties: The room temperature ultimate tensile strength (UTS) of the investigated hybrid composites is shown in Fig. 7. For a fixed values of graphite content and squeeze pressure, increasing SiCp improves considerably UTS. The relative increase is quite remarkable for composites squeezed at lower pressure (50 MPa) and having higher graphite content (5%). Moreover, the increase in UTS with SiCp is not absolute since the data start to decrease for composites having 15% SiCp. On the contrary, graphite addition seems to have an adverse effect on UTS. This is quite true also for 0.2% proof stress (Fig. 8) It is obvious that the graphite particles do not contribute to the strength properties of the composites. The SEM examination of polished surfaces beneath the fracture surfaces indicated cracking of particle-matrix interface. The particle matrix interface thus undergoes separation under tensile loading and void formation takes place on the particulate phase. The decrease in strength properties with increasing wt'/f- of -34- 350 350 300!- 250 o> 200 150 100 -*-&* - o •*-Gr% - 3 ••• Gr% - 5 50 1 5 10 15 . %SICp XSiCp Fig.7a:Varlation of ultimata (ensile strength with SiCp-contenl for various wt.% graphite under: Fig.7b:Varialion ol ultimate tensile sirenglh with a) SO MPa- squeeze pressure. SiCp-contenl for various vvt.% graphite under: b) 125 MPa- squeeze pressure. 280 240 240 ,200 200 a. • 7 / S 160 160 / / 120 . / 1- / 6 ' 60 6'i ~Gr% -0 40 = 40 "" 3r% - 3 m * Or% - S • -Or%-t n 1 i, , 10 15 10 15 %SiCp KSiCp Flg.Sa:Varlallon of 2% Proof slresi with SrCp-conlenl Fig.eb:Variatlon of 2% Proof streas with SiCp-content for various wt.% graphite under: for various wt.% graphite under: 50 MPa- squeeze pressure. 125 MPa-squeeze pressure. -35- graphite is apparently due to the fact that the graphite particles act as voids within the continuous eutectic phase, reducing the load-bearing area of the tensile specimen and therefore the breaking load. Additionally, the stress acting at the graphite-matrix interface may cause crack nucleation to occur at lower magnitudes of stress. However, the negative effect of graphite on strength properties can be neutralised through the application of pressure during solidification (i-e squeeze Casting) and addition of SiCp. Squeezing leads to an intimate contact and efficient bonding between the matrix phase and the reinforcement phase with an effective laod transfer from the matrix to the hard phase. It has been recognised that these can be a problem associated with the production of MMCs using liquid metal production routes in view of the possibility of retaining remanant porosity between the matrix and the reinforcing phases. The decrease in UTS for composites having high SiCp - content (15%) can be attributed to the high density of imperfections (voids, microcracks ...) associated with high contents of SiCp- such decrease in UTS is quite observed at lower squeeze pressure (50 MPa) rather than at higher values (125 MPa). This signifies further the relative importance of squeeze casting technique. On the other hand on decrease in 0.2 % prof stress was observed with in the examined range of SiCp as shown in Fig. (5). So, one can say that the imperfections associated with SiCp additions have no significant effect on proof stress. The variation of ultimate tensile strength of the matrix material (Al-12% Si) and hybrid composite material with test temperature and squeeze pressure is shown in Figs. 9) and 10 respectively. The grain refining effect coupled with microstructural homogenity associated with stirring and squeeze casting the considerable improvement in high temperature UTS compared with conventional casting (MPa) of the matrix material (Fig. 9). -36- The UTS of the hybrid composite falls less rapidly with increasing test temperature as shown in (Fig. 10) and it remains significantly higher than that of the base alloy (Fig. 9). The effect of SiCp on the deformation is expected to change as the temperature is raised. At room temperature, dislocations accumulate at the particles during deformation, and this can provide both nucleation sites and a large driving force for particle-stimulated nucleation of recrystallization on subsequent annealing' '. However, during deformation at elevated temperature, dislocations may be able to climb around the particles, and thus no accumulation of dislocations can occur. In this case, no deformation zones will be formed and the dislocation structure is expected to be rather similar to that of unreinforced alloy1151. The dependence of elongation % (at room temperature) on SiCp-content and squeeze pressure of Al-12 Si/3 Gr composites is represented in Fig. 11. For a particular pressure, increasing silicon carbide particles reduces the ductility, simply due to the high hardness level of the ceramic phase. The reduced elongation caused by reinforcement is reported by many researchers^ '. Meanwhile, the positive effect of stirring and squeezing on improving ductility through grain refining as well as better homogenity is observed from Fig. 11. The Rockwell hardness of Al-12 Si/3Gr composites squeezed at 80 MPa before ad after T6 heat treatment (ST at 520*C, 8h quenched in warm water at 60°C then aged at 180°C, 6h) is shown in Fig. 12. The increase in hardness with SiCp - addition and after T6 thermal treatment results from the high dislocation density in these materials and the higher hardness of SiCp. The high dislocation density is generated during cooling from the annealing temperatures a result of the large difference in coefficient of thermal expansion between Al and SiCp. This may also explain the increase in hardness with increasing SiCp content. -37- Resistivity Measurements Fig. 13 shows how the electrical resistivity, measured at room temperature, was affected by SiCp and graphite additions. Increasing SiCp-contents results in an increase in resistivity for all the examined ranges of material as well as processing variables. Such increase can be attributed to the high density of imperfections associated with SiCp (voids and microcracks). High values of graphite-content and squeeze forming pressure are shown to increase the resistivity since graphite particles act as voids. Setn Proctography Examination of the different fracture surfaces in the SEM shows that the extent of ductile dimpling is decreased progressively with increasing SiCp and graphite particles (Fig. 14). Matrix cracks were invariably noticed especially in the case of higher wt% of graphite. Also, the graphite particles were pulled out from the matrix leaving behind cavities. The separation at the particle-interface and the matrix cracks in the vicinity of particles is believed to be instrumental in decreasing the strength properties of the composites. In the composites material, fine dimples occur in the regions of matrix between the particles which have undergone ductile tearing (Fig. 15). From these observations it is apparent that the fracture process in these hybrid composites is quite complex, involving several different phenomena. Voids associated with clusters of particles will open up and grw during straining. It is not clear whether the failure observed in the particle-matrix interface is due to the opening up of incipient voids or to intrinsic failure of the interface. In addition to voiding, graphite and some SiC particles, and any coarse intermetallics associated with them, crack during tensile straining. Cracks produced by voiding or by particle fracture then link by ductile failure of the matrix. The component illustrated in Fig. 16 has been made -38- 6 ?W>- \ -v V S 5 'SO * r X * I too § so A -I I —I I — ..I _.•_.! I. I - 0._ i .„• . i • . . i .. i .. . ~0 SO «00 I40 20O J50 300 35O40O4MSO0 0 50 lOO ISO ?00 ?W *on?f»0 400 4SO *•• T»«t l»"*ti«»irtu«« "C Fig O Yh» ri«n«*Mj«fc« nl utllmnl* <*M«II« «4umptlt n< !**» maltiR aMoy (AM 3 Si) on !••! l«mp««*tur« (AM 2 S./10 S«Cp * 3 10 MP» i •o Mr» mur< /K ID a w / c a x: *9 J* »*« 16 heal t»>atnie i... 5 10 1S 10 \.RICp % SIGH Fig 17 nocVwrtl hnrdn*»«*«» U b«*ot* «nd »M«»t lr h*;it Fig 11. Ell«cl ol SICo-«ddlllont on lh» nlongatinn lr«atin«nt ol AI-12 SI/31!' eomnosim at allarlnO % ol AM 2 Si/3Gr eomixMltn satMWtnd undar diWemnt hy SiCn ndrtnions. tH» MP« s««in«j* pi nsmirn fjf«tiuf«s. i»-.-rtnr:..-|tic i ciii lflt».aL«!!il£!'f»'."» J |Tii5'::ct!.,'"a'""'1 SU.iin, ll,.~a.H la i i««t« wtm«a Into whlchf |-i'-.r.nJuu-.i;sin^j»7,.i»,.-?ii;fr • LKSSVAJHU J*-1iilJ-0'!Ct.-.. J Mil I [M., l.[„lr.TI |K>ifinliMT»>:i 111.. 11.. .'.'> |r..ri.-*4«l l.i;« E O . _ - - M 1=0 8 " _ y' * *• •**• Ml. O S 10 15 *.SlCp Flo 13' E'ectrlcnl reilatrvlly at room temperature as affected by SICp- and graphite eddlllons and •queue preaiure Fi.0 14 SE* free I adraphu of the aurfaca of twil I • Called coapeetteat al *i-ia si, no HP*. ion* bl 41- 12 SI /S SICpOCr. SO Hra.lOOl c) *l- 12 SI /10 SIC|-X.r. 80 Ufa. 1 00« d) Al- 12 SI /IS SlCrX.r, 10 MPa, I OOx -40- FJLg 15 5F:tt Mnrl(i Fiq 16 A" »rti<:l.o made from Al- 12 Si/10 SiCp- 3Gr, 110 MPa -41- from AI-l2oi (10 SiCp + 3Gr) hybrid composites squeeze formed under 80 MPa. This component is used in the clutch assembly of fiat 128 to replace another made from steel which had faced several problems. With the present mater'.al and processing, there is not such a great need for a reasonable fluidity in order to reproduce the die features faithfully. The component represented in Pig. 16 compete both technically and economically with alternativenes made by low pressure and gravity die costing, high-pressure die casting, forging and fabrication. Concluding Remarks 1- The combination of vortex technique with squeeze casting have succeeded in producing hybrid composites having a uniform distribution of the particles, near-net shape, mechanically round and finer structure. 2- The introduction of graphite particles offers the possibility of producing Al-Si/SiCp composites with a good property profile Aluminium hybrid particulate composites were characterized by lower cost and more isotropic properties compared with fiber reinforced composites. Excessive chemical reaction between SiCp and the reactive aluminium melt, that degrades strength properties, can be avoided. 3- The mechanical properties measured suggest that these materials provide a reasonably high strength (about 280 MPa) at appropriate ductility (7%), compared with the monolithic alloy. 4- Void formation mechanism appeared to play a dominant role in the failure of the investigated composite tensile tested at room as well as high temperature. 5- The process outlined in this paper for forming hybrid composites may have general applicability to other systems. Our future work will be focused on producing a piston from these materials. -42- 6- It is recommended for a failure work ot investigate the tribological behaviour of the present hybrid composites as affected by: materials, processing and test variables. Also, it might be fruitful to study the aging response at 180'C (T6 treatment) and at room temperature (T4 treatment) to quantify the role of particulate additions. In this regard, resistivity measurement seem to be a very sensitive method to follow the structural changes. References 1. Das, S., Prosed S.V. and Ra mac hand ran, T.R. "Microstructure and wear of cast (Al-Si Alloy) graphite composite", wear, 133 pp 1731187 (1989). 2. Khedre, O.M., Attia, A. Gado, M., El-Desouky, A. and Kafagy, A. "Damping characteristics and wear behaviour of sintered All Graphite composites" key Eng. Materials, vol. 79-80 PP 349/364 (1993). 3. Japanese patent No. 85 054815. 4. Krishnan, B.P. und Rohatgi, P.K. "Modification of Al-Si Alloy Melts containing graphite particles dispersion, Metals Technology, vol. 11 pp. 41-44, Feb. (1984). 5. Surappa, M. and Rohatgi, P. "Production of Al-Gr particles composites using copper-coated graphite" Metals Technology PP 358/361 Oct. (1978). 6. Gbrbunov. V., Parshin, V. and Panin, V. Russ cast prod., PP. 348/349, (1974). 7. Teikoku Piston Ring Co. Ltd: Japanese patent 4800 311, (1982). 8. Sritvatsan, T.S. Ibrahim, LA., Mohamed, F.A., and Lavernia, E.J. "Processing techniques for particulate - reinforced metal Al matrix composites, J. Mat. Sci, 26 PP. 5965/5978, (1991). 9. Rohatgi, P.,m "Cast Al matrix composites for automotive applications", JOM, PP 10/15 April (1991). 10. Srefanescu, DM. "Issues in liquid processing of particulate metal matrix composites," key Eng. Materials, Vol. 79-80 PP. 75/90, (1993). -43- 11. Geigcr, A.L. and Jackson, M., "Low expansion MMCs boost avionics" Advanced Material and Processes, 7PP 23/30(1989). 12. Comic, J., Mortensen, A. Flemings, M., in S.L. Mattews et. al (eds) Proc. 6th Int. Conf. of Composite Materials, London, Vol. 2, Elsevier, Amsterdam. 2297/2319 (1987). 13. Modolfo, L.F.: "Aluminium Alloys: Structure and properties", Butterworths (1974). 14. Humphreys, F.J., in Loretto, M. (ed), Dislocations and properties of real materials. Institute for Metals, London 175 (1985). 15. Humphreys, F.J. "The thermomechanical processing of Al-SiC particulate composites" Materials Sci. and Eng., A 175 2671273 (1991). 16. Verma, S.K. and Dorcic, J.L.: "Manufacturing of composites by squeeze casting", in S.G. Fishman and A.K. Dhingra Oeds), Cast Reinforced Metal Composites, ASM int. Metals Park, O.H. (1988). •44. £ C f£ O slSolf First International Spring School & Symposium on Advances in Materials Science (SAMS 94) 15-20 March 1994 IMPACT AND FRACTURE TOUGHNESS PROPERTIES OF AUSTEMPERED DUCTILE IRON F.S.Ahmed and M.M. Ghoncim 1. Shoubra Fac. of Eng.,Zagazig Univ., Cairo, Egypt 2. Metallurgy DepL,Atomic Energy Authority, Cairo, Egypt ABSTRACT Fracture characteristics of bainitic ductile iron transformed at various austem- pering temperatures and austempering times are evaluated by plane strain fracture toughness test and impact test Impact tests were carried out on standard V-notch Charpy specimens at temperature range -100 to 350 °C. Fracture toughness was measured using the short rod (SR) test procedure. Standard compact tension (CT) specimens of 25 mm thickness were also tested for reasons of comparison with the SR results. Optical micros copy and scanning electron microscopy (SEM) were used for the examination of micros- tructure and fracture surface morphology, respectively. Austempering at 350 °C resulted in the optimum impact energy and fracture toughness properties, whereas that at 270°C showed the inferior properties. Fractographic exanunation showed that different fracture mechanisms are associated with the different impact and fracture toughness levels obtained. 1. INTRODUCTION The production of austempered ductile iron (ADI) involves two processes; metal casting and heat treating. These processes are closely related and must be coordinated for a successful application [1-3]. Casting of an ADI part must be adjusted with respect to chemical analysis and castirg technology. Optimum mechanical properties are obtained when the graphite structure is fully nodular. The mechanical properties of DI are significantly enhanced by austempering heat treatment. The process involves heating to the austenite temperature range until complete austenitization is achieved followed by quenching and holding in a bath of specified temperature to obtain bainitic structure. Both austempcring temperature and holding time at the bath temperature have strong effects on the microstructure and mechanical properties [1-7]. Austempered ductile irons (ADIs) possess a fairly attractive combination of mech anical properties; high strength with acceptable ductility, toughness and wear resistance. They find increasing use in many applications such as crank shafts, gears and transport containers of nuclear waste. Reductions of 30% or more in component cost typically result when DI is substituted for cast steel due to lower melting point and better machinablity [8], However, DI is still excluded from many applications because of the concern over its presumed low fracture toughness indicated by Charpy impact values which seldom exceed 20 J. However, recent publications [81 reported values of fracture toughness of about 80 MPa m for AD Is, which is high enough for many applications. Nevertheless, the published fracture toughness data of ADI is scarce and more work is still required in this field. Fracture toughness testing is relatively difficult and expensive [9,10]. The recently proposed short rod (SR) method for measuring fracture toughness [9-11J appears to offer a simpler and less expensive method compared to the generally accepted ASTM E399 standard method, especially for materials of high strength. Short rod (or short bar) testing method, presented by committee E24 of ASTM, uses a simplified specimen configuration to determine Kic value. The specimen shape is either cylindrical or rectangular with a deep chevron notch, Fig. 1. During testing both the load and crack mouth displacement are recorded. This test docs not require previous precracking [9] because the specimen geometry is such that a crack having critical length "a" is created when the maximum load "Pc" is reached, after which the crack grows in an unstable manner. Provided that conditions of linear elastic fracture mechanics (LEFM) apply, Fig.2, the fracture toughness Kic is given by m Kfc =APC lD (1) where: A - short rod calibration constant = 22 P0 = maximum load D = Short rod diameter Barker [9] has modified Eq. I to account for localized crack tip plasticity "p". Crack tip plasticity appears as non linearity in the load-dcflcction record, as indicated in Fig.2, and is determined from the loading and unloading cycles, defined as the increase in residual mouth opening between two relaxations divided by the increase in mouth opening at the average crack-advancing load between the two relaxations, p = A Xa/ A x. In this case Kic is given by 3/2 ,/2 Kfc =(APc/D )((l+p)/(l-p)) (2) In the present study the influence of the DI microstructures obtained after various austempcring treatments on the Charpy impact and fracture toughness properties was investigated. The results were discussed in light of the fracture surface morphology examined using scanning electron microscopy. -46- 2. EXPERIMENTAL 2.1. Material preparation and heat treatment Ductile iron Y-blocks of 30 mm thickness were cast vertically in sodium silicate CO2 sand mould,in a manner reported previously! 12], The average chemical composition wt% is: 3.51 C, 2.4 Si, 0.23 Mn, 0.013 P, 0.01 S, 0.046 Mg, 0.47 Cu. the remainder is Fc. The produced DI has 250-300 nodulcs/mm2 and nodularity above 90%. Short rod (SR), Charpy V-notch and compact tension (CT) specimens were machined from the test section of the Y-blocks. The diameter of the SR specimen, Fig. I, was 26 mm and the thickness of the CT specimen, Fig.3, was 25 mm. The specimens were then subjected to austempcring heat treatment by austcnitization at 980 °C for one hour, followed by quenching in salt bathes at temperatures 270, 350 and 400 °C and isothermal holding at temperature for 5,15,30,45 and 60 min followed by oil quenching to room temperature. The austempcring time for the impact specimens was 60 min for the three austempcring temperatures. Some of the SR specimens austempcred at 270 C for 5 to 60 min as well as some of those austempcred at 350 and 400 C for 5 min were further tempered at 200 C for 3 h. 2.2. Testing Impact tests were performed using an Amslcr impact machine. Short rod and compact tension specimens testing was carried out on a Schcnck scrvohydraulic universal machine. Fatigue precrackingof theCT specimens was performed on this machine also. CT specimens were fatigue precracked and tested in accordance with ASTM E399 standard method [13]. Precracking was carried out so that the ratio of the crack length to specimen width was about 0.5. Both the SR and CT specimens were tested under displa cement control of 0.5 mm/min. Microstructure was examined using optical metallography while X-ray diffraction was used to measure the amount of retained austcnitc. Fracture surfaces of impact and SR specimens were examined using scanning electron microscopy. 3. RESULTS AND DISCUSSION 3.1. Microstructure Figure 4 shows the microstructure of the DI in the as cast condition as well as alter austempcring for 60 min at 270, 350 and 400 °C. The variation of the amount of retained austcnitc with lime for the three austempcring temperatures is shown in Fig.5. These changes in microstmcturc arc in agreement with the general behaviour of DI throughout the austempcring process [3J. «47- Fig.l: Short rod (SR) specimen ——-1 i fT 50.80 -4 •j. **W b r * •• ."^5" - Fig.2 : Schematic for the load-displacement Fig.3 : Standard ASTM E399 compact of the short rod specimen testing tension (CT) specimen -48- Fig.4 : Optical micrographs of DI and ADI: a. as cast b. austcmpered at 270 °C, 60 min c. austempetcd at 350 °Cr 60 min d. austempcrcd at 400 °C, 60 min -49- The purpose of auslcmpcring DIs is to produce a matrix structure consisting of bainitcwilh high-carbon stable austenite (retained austenite). As the austempcring temperature increased from 270 to 400 °C, the transformation products changed from primarily lower bainite (fine structure), Fig.4-b, to upper bainite (coarse structure), Fig.4-d. Austempcring at 3S0 °C resulted in a mixture of lower and upper bainite. The amount of retained austenite increases with transformation temperature, attaining values of 18,32 and 45% for austempcring temperatures 270, 3S0 and 400 °C , respectively . For each austentpering temperature, the amount of retained austenite increases with time. At short times small part of the austenite is transformed to bainite and the amount of rejected carbon is consequently small. Therefore, the carbon content of the unreacted austenite is lower than that necessary to stabilize the austenite, which is transformed to martensite on cooling to room temperature. With increasing the holding time both the amount of bainile and that of the carbon-enriched untiansformcd austcnitc (retained austenite) increase and the amount of martensite is reduced. 3.2. Impact and fracture toughness properties The variation of the Charpy V-notch impact energy with test temperature for the 270,350 and 400 °C austempering temperatures is shown in Fig.6. Austempering at 270 °C nearly did not affect the impact toughness all over the testing temperature range. However, pronounced improvement took place after austempering at 350 and 400 °C. Fractographic examination results of Charpy specimens are given in Fig. 7. Fracture toughness values obtained from the SR specimens according to equation 1 are given in Table 1. All the specimens showed linear elastic behaviour. On the other hand, the results of CT specimens testing are given in Table 2. The conditions for linear elastic fracture in this case are : 2 Pmal?Qil.l,miB,ii 2.5(K where Pm« = the maximum load, PQ = the load (conditional load) used to calculate KQ , KQ = conditional fracture toughness; KQ - Kb if the above conditions are fulfilled, B = specimen thickness, a =• crack length, and cry » yield strength. The ratio Pm« /PQ was less than 1.1 for all the specimens. The yield strength for similar conditions [7] is about 800 MPa. Using the values given in Table 2, the smallest thickness B or crack length a is 9-17 mm which is smaller than those used in the present study ( a = B = 25 nun). Therefore the results are considered valid linear elastic fracture toughness Kic. Tables 1 & 2 show that fracture toughness values obtained by SR tests are higher than the corresponding values of the CT specimens. It is believed that this was due to the -50- 9. L.M.Barkcrand FJ.Baratta, J.TcstEval., 1980, pp. 97-102. 10. R.W.Hcrtzberg, Deformation and Fracture of Engineering Materials, Second edition .John Wiley & Sons, New York, 1989. 11. J.Arnazabalelal,Mater.Sci.Tech.,Vol8,1992, pp.263-273. 12. F.S.AJuncd.The Egyptian Foundrymen's Soc.Mag.,1992,pp.27-39. 13. ASTME399-78a, 1980 Book of ASTM Standards, Part 10,1980. 14. L.EIdoky and R.C.Voigt,AFS Trans. Vol.93,1985,pp.365-372. 15. C.Bak, M.Degois and J.M.SchissIcr,AFS Trans.Vol.88, 1980, pp.301-312. 16. RN. Wright and TRJarrell, AFS Trans. Vol.93,1985, pp.853-866. 17. T.Luycndijk and H.Nieswaag,Proc. 50th Int.Foundry Congress, Cairo, 6-11 November 1983, Paper No.9. 18. T.S.Shihet at, AFS Trans. Vol.99,1991,pp. 793-808. Table 1: ADI fracture toughness values from SR specimens testing, equation I; A = 22. austempering fracture toughness, MPa m temperature, °C holding time,min 5 IS 30 45 60 270 13 21 34 50 55 350 40 57 - 75 80 400 21 57 59 68 70 Table 2: ADI fracture toughness values from CT specimens testing austempering fracture toughness .MPam"2 temperature, °C holding time, min 15 60 350 49 67 400 50 S2 -51- 10 20 30 40 SO Austempering Time, min Fig.5 : Effect of austempering time on retained austenite content for various austempering temperatures J AC : as cast • A • austemper cd, 2701., Go min a 8 : . 350t. 00 min • C : . UXit. 60 min 20 I- • • c • • 16 - /" a / • 1 0 0 B 0 o c 3 UJ j/6 »/ ' A •© •0 o - 8 . / , • U ^^-o- o^ • • a. ^AC E ~ A i . 1 1 1 ..1 1 1 1 •100 0 100 200 300 400 Test Temperature, *C Fig.6 : Effect of test temperature on the impact energy of the as cast DI and AD I. -52- value of the constant A in equation 1. The value A = 22 was derived for a SR specimen with initial crack length a = 0.45 D, where "D" is the specimen diameter [9], i.e., for a = 12 mm in the present case. However,"a" in the present investigation was about 8mm. Accordingly, the constant A must be adjusted. This can be done using the values of the fracture toughness obtained by CT specimens for calibration, which gives an average value of A =18. This value is used in Fig.8 which shows the variation of the fracture toughness of the ADI with holding time. SEM fractographs of SR specimens are shown in Fig.9. 3.2.1. Effect of austempering temperature The above results showed that the impact energy and the fracture toughness properties of the investigated DI are strongly dependent on the austempcring conditions. Austempering at 350 °C gave the optimum impact energy and fracture toughness at room temperature. On the other hand, austempering at 270 °C showed the poorest results and the material was nearly as brittle as the as cast one. This is in agree ment with the results obtained by other investigators [4,14-17]. Lower bainite with small amount of retained austenite, obtained at 270 °C in the present case, is characterized by very high tensile strength and low ductility. The microstructure produced at 400 °C is upper bainite with high retained austenite content. This microstructure is known to have a relatively lower strength but with considerably higher ductility [1-3,12,18]. The mixed structure, obtained at 350 °C, possesses the optimum combination of strength and ductility and, consequently, superior impact energy and fracture toughness properties. The DI austempered at 400 °C for 60 min showed a rapid increase in impact energy at testing temperatures > 50 °C to attain the highest upper shelf energy. This could be attributed to the effect of the retained austenite (45%) in these specimens, Fig.5. It was shown [11] that ADI of microstmctures with high retained austenite content, produced at about 400 "C, have higher value of strain hardening compared to those with low retained austenite content, produced at 300 °C or lower. The material transformed at intermediate temperatures (e.g., 350 °C) has intermediate strain hardening value. As the test temperat ure increases the chance for plastic deformation increases and, hence, the effect of higher strain hardening, due to the higher retained austenite content, becomes more evident Retained austenite compensates for the softening due to graphite nodules debon- ding which takes place at the upper shelf energy range. SEM fractographs of impact specimens austempercd for 60 min at 270, 350 and 400 °C are shown in Fig.7. Fig.7-a shows that the fracture mechanism in the as cast specimens was transgranular cleavage. This behaviour took place at room temperature as well as at elevated temperatures. Austempered specimens tested at room temperature showed predominantly ductile dimple fracture with some transgranular quasicleavage features, Figs.7-c & 7-d. The quasicleavage proportion increases slowly with decreasing test temperature. At -IO0°C, Fig.7-b, the fracture was completely by quasicleavage mode. Above room temperature the fracture was completely ductile dimple. Fig.7-e is for a specimen austempered at 350 CC and tested at 100 °C (E = 14 J), while Fig.7-f is for a specimen austempcrcd at 400 °C and tested at 200 °C (E = 20 J). These two fractographs -53- ••• n vVu» • Fig.7 : SEM fractographs of Chaipy V-notch impact specimens of DI and ADI (continue) -F4- Fig.7: SEM ftacographs of Chaipy V-notch impact specimens a. as cast, tested at 1W#C b. austempered at 400°C/60 min, tested at -100°C c. austempered at 270°C/6O min, tested at 25°C d. austempered at 400°C/6O min, tested at 25°C e. austempered at 350°C/60 min, tested at 100°C f. austempered at 400°C/60 min, tested at 200°C show dimple coalescence and dcbondcd graphite nodules with high local plastic deform ation. The microplasticity around the dcbondcd nodules is much higher than that at room temperature, while the brittle quasiclcavagc case occurs with little or no microplasticity around the nodules or in the matrix. The quantity of nodules observed in the fracture surface in the case of ductile dimple fracture is greater than that for cleavage or quasi- clcavr.jc fracture. The observation that ductile dimple mode was predominant even at subzero test temperatures is very important. This represents an advantage for the material applicati ons in machine components subjected to impact loading at low temperatures. 3.2.2. Effect of austempering time Figure 8 shows that fracture toughness increases with increasing the austemp ering time for the three austempcring temperatures studied. This is attributed to the corresponding changes in the microstructure. Short treatment times produces high amounts of martensitc in the final structure. Due to the presence of this brittle phase in the material the fracture propagates mostly intergranularly through the prior austenite grains resulting in a very low fracture toughness. Figs.9-a & 9-b show the intcrgranular mode of fracture observed at short austempering times. These fractographs show that the quasiclcavagc mode, which is tougher than intergranular mode, was also operating but to a much less extent. Increasing the austempering time results in decreasing the martensite content in the material. This causes the fracture toughness to increase; the intcrgranular fracture proportion decreases, the quasicleavage proportion increases and the possibility for dimple fracture increases. Figs.9-c & 9-d present the fracture surface of SR specimen with relatively long austempcring time. In this case the fracture took place through a mixture of ductile dimple, quasicleavage and intergranular mechanisms. 3.2.3. Effect of tempering Figure 8 shows also the effect of tempering at 200 °C for 3 h on fracture toughness of austempercd SR specimens. This treatment produced an appreciable increase in fract ure toughness in the case of specimens austempered for short time (5 min). However, fracture toughness of specimens austempered for longer times was slightly reduced. Several investigators [14,15] indicated that tempering of AD I structures causes the form ation of secondary graphite. The nucleation of this secondary graphite depends on the prior matrix structure. Tempering of martensite-containing structures results in a fine distribution of secondary graphite throughout the matrix leading to improvement of toug hness. Conversely, tempering of bainitic structures results in the formation of secondary graphite rings around the primary nodules which causes a reduction in the fracture toughness [14,15] -56- • Justcmperto t\ 270 C .. 3r0 °C - «» °C 80 \ &••'» tempered at ?0D°C/3h alter no a. austempering 2 60 1/1 S c x .8 40 2 20 J_ 10 20 30 40 50 60 Austempering Time, min Fig.8 : Effect of austempering time on fracture toughness of ADI -57- Fig. 9: SEM ftactographs of SR specimens of ADI : a,b: specimens austempered at 270 *C tor 15 min c,d: specimens austempered at 350 °C for 45 nun -58- 4. CONCLUSIONS In the present investigation a pcarlitic cast ductile iron was auslcmpcrcd at 270, 350 and 400 °C for times from 5 to 60 min. The main conclusions arc : 1. Austcmpcring at 270 °C produces lower bainitc while that at 400 "C produces upper bainitc. The microstructurc aficr austcmpcring at 350 °C contains a mixture of lower and upper bainitc. 2. The amount of retained austcnitc increases with increasing the austcmpcring temperature and/or the austcmpcring time. The maximum values arc 18,32 and 45% after austempcring for 60 min at 270,350 and 400 °C, respectively. 3. Austempcring at 350°C for 60 min gives the highest value for both impact energy and fracture toughness whereas the lower values arc associated with austcmpcring at 270 °C. 4. Fracture toughness increases with increasing the austcmpcring time for the three austempcring temperatures. This can be attributed to the corresponding increase in the retained austenite and reduction of martensite in the microstructurc. 5. The variation of impact energy and fracture toughness values is related to the variation in fracture mechanisms. Lower values correspond to cleavage fracture in the case of as cast DI, and to intcrgranular and quasiclcavagc in the case of ADI. The higher values correspond to the presence of transgranular ductile dimple fracture mode. Dcbonding of nodules and local plastic deformation around them increases with the increase in fracture toughness. 6. Tempering at 200 C for 3 h, after the austempcring treatment, improves the fracture toughness in the case of short austempering time (5 min), while slight reduction takes place for longer austempering times. REFERENCES 1. B.V.Kovacs, AFS Trans. Vol.99,1991, pp.281-286. 2. D.J.Moore, R.T.Rouns and K.B.Rundman, AFS Trans. Vol.95, 1986, pp.765-774. 3. R.A.Hardingand G.N.GiIbcrt,Br.Fndry.Vol79,1986,pp.489-495. 4. RK.Nanstad, J.Worzala arid C.R.Loper,AFS Trans. Vol.83, 1975,pp245-256. 5. S.A.El-Dalil and F.S.Ahmed, Modelling, Simulation and Control B AMSE Press, Vol.32,1990,, pp.51-64. 6. F.S.Ahmed, and A.A.El-Zoghby, Egyptian Foundrymcn's Society Magazine", 1988, pp. 34^*5. 7. J.L.Doong and C.S.Chen, Fatigue Fract. Engng Mater.Struct. Voll2, 1989, pp. 155-165. 8. W.L.Bradley, K.E.Mc Kinney and P.C.Gerhardt, ASTM STP 905, 1986, pp.75-94. EG fS>o/ISo 5 First International Spring School & Symposium on Advances in Materials Science (SAMS 94) 15-20 March 1994 INSTRUMENTED IMPACT PROPERTIES OF SOME ADVANCED NUCLEAR REACTOR PRESSURE VESSEL STEELS M.M.Gnoncim,A.M.NasrcIojin,A.A.EIsayed,D.Pachiir an(j F.H.Hammad Atomic Energy Authority, Cairo, Egypt • KFA, Juelich, Germany ABSTRACT The steels used for the construction of nuclear reactor pressure vessels are low alloy ferrilic steels. These steels must have good impact properties, i.e., low transition temperature and high upper shelf energy, both before and during being subject to service conditions', the most important is neutron irradiation. Extensive R&D work has been conducted to make the production of such steels possible. In the present study instrumented impact testing was conducted for three advanced pressure vessel steels in comparison with a conventional pressure vessel steel. The first one of the advanced steels was a 20MnMoNi55 (ASTM A533-B C1.2) weld produced in Germany (GW). The second was an ASTM AS08 C1.3 forging produced in France (FF). The third was an ASTM AS33-B Cl.l plate produced in Japan (JP). The conventional steel was an ASTM AS33-B C1.1 plate (HSST 03 plate) produced in USA(HSST) and it represents steels used in most nuclear reactor pressure vessels currently in operation. Both microst- ructure and fracture surface were examined using optical and SEM techniques. In general, the advanced steels showed much better impact properties (lower ductile- brittle transition temperature and higher upper shelf energy) than the conventional steel. The ductile-brittle transition temperature (DBTT) at 41 J (T41J) was -73,-57, and -44 °C for the GW, FF, and JP steels, respectively, compared to 6 °C for the HSST steel. The advanced steels showed an upper shelf energy of 183-200 J compared to 125 J for the HSST steel. The load-time traces showed that the increase in the fracture energy was mainly due to the increase in the fracture propagation energy rather than the initiation energy. The initiation energy of the HSST steel was about 50 J while that of the advanced steels was about 60 J. On the other hand, the propagation energy for the HSST steel was 75 J compared to about 130 J for the advanced steels. The improvement in the toughness level of the advanced steels compared to that of the HSST steel was related to the difference in chemical composition, microstructure and fracture surface morphology. -61- KEYWORDS Pressure vessel steels; impact toughness; baiting microstructurc; fraclography. 1. INTRODUCTION More than three quarters of the nuclear power stations all over the world already in operation or under construction use steel pressure vessels to house the reactor core. Successful and safe performance of the power stations depends on the reliability of steel pressure vessels [1J. The ability of the reactor pressure vessel to resist brittle fracture is of particular importance. This property depends on the material fracture toughness. In general, the steel fracture toughness decreases, and therefore the risk of brittle fracture can be considerably enhanced, due to neutron irradiation during reactor operation. Impact Charpy testing is used to determine the fracture toughness of pressure vessel steels (PVSs). The two important parameters defined by this test arc the ductile-brittle transition temperature (DBTT) and the upper shelf energy (USE). Irradiation shifts the DBTT to higher temperatures and lowers the USE . The magnitude of this effect (irradiation sensitivity) varies from one steel to another depending on several factors [2]. The most important is the chemical composition. Residual elements especially Cu and P were found to increase irradi ation embrittlement, with Cu being the most serious residual clement . It is therefore important for the pressure vessel steel to have: (i) high initial fracture toughness (low DBTT and high USE), (ii) low irradiation sensitivity, mainly through minimizing the residual element content. In addition, new designs of the nuclear pressure vessels adopt using of more forgings [3,4]. This minimizes welds and especially eliminates longitudinal welds, which facilitates in-service inspection. Furthermore, forging produces more uniform material and the directionality of the mechanical properties is reduced compared to steel plates. The present work is concerned with studying the initial impact properties of three adva nced pressure vessel steels in comparison with a conventional pressure vessel steel as a reference material. 2. EXPERIMENTAL The materials used in the present investigation were three steels representing the advanced PVSs as well as a reference steel representing the conventional PVSs [lJ.The first one of the advanced steels was a 20MnMoNi55 (ASTM A533-B C1.2) weld produced in Germany (GW). The second was an ASTM A508 C1.3 forging produced in France (FF)- The third was an ASTM A533-B CI. 1 plate produced in Japan (JP). The conventional steel was an ASTM A533-B Cl.l plate (HSST 03 plate) produced in USA (HSST) and it repres ents steels used in most nuclear reactor pressure vessels currently in operation. The Chemical .62- composition and heat treatment arc given in Tables 1 and 2. Standard Charpy V-notch specimens 55x10x10 mm with notch depth 2mm were machined from the four steels . The specimens were cut with their axis parallel to surface, perpendicular to rolling direction (plate), in tangential direction (forging) and the notch perpendicular to surface. Impact testing was carried out according to ASTM E23 [5]. An instrumented impact testing machine with a total energy 300 J was used. Impact tests were conducted over a temperature range to generate full transition curves. The load-time traces produced from the test were utilized to obtain the dynamic yield strength and the fracture initiation and propagation energy. Optical micrographic observation was conducted by etching polished specimens using a solution of 4% picric acid and 1% nitric acid in metha nol. Fracture surface morphology was examined using scanning electron microscopy. 3. RESULTS 3.1. Microstructure The microstructure of the investigated steels is presented in Fig. I. The HSST and the FF steels showed tempered bainitic structure. The JP steel showed a mixture of procutectoid ferritic and bainitic structure whereas the GW steel showed an acicular ferritic structure. The GW steel was found to have the finest grain size, about 5 |im, while the FF steel had the coarsest grain size, about 30 ftm. The JP and the HSST steels had grain size of 15 and 20}im, respectively. 3.2. Impact Properties The ductile-to-brittle transition curves of the tested steels are shown in Fig.2 In general, the advanced steels show much better impact properties (lower ductile-brittle transition temperature and higher upper shelf energy) over the conventional steel. The ductile-brittle transition temperature (DBTT) at 41J (T41J) and the upper shelf energy (USE) values are given in Table 3. The GW steel showed the lowest T41J (-73 °C) while the HSST steel had the highest value (6 °C). The advanced steels showed an USE of 183-200 J compared to 125 J for the HSST steel. Load-time curves for the tested steels were obtained at different temperatures. Those representing specimens tested at room temperature are shown in Fig.3. The GW and JP steels show upper shelf behaviour. The FF steel was in the upper part of the transition region while the HSST steel was in the lower part of the transition region. The load time curves were utilized to obtain the dynamic yield strength (ayt) [6] and the fracture initiation energy (Ej) and propagation energy (Ep). c^ values are presented against the testing temperature in Fig.4. The GW and FF steels show cyd values higher than those of the JP and HSST steels. Table 4 presents values of the dynamic yield and dynamic ultimate -63- Table I: Chemical composition, wt.% steel code GW FF JP HSST c 0.08 0.15 0.18 0.25 Si 0.17 0.28 0.22 0.25 Mn 1.45 1.3 1.4 1.33 Mo 0.61 0.5 0.58 0.51 Ni 0.93 0.7 0.66 0.65. Cr 0.02 0.24 0.2 0.1 P 0.011 0.009 0.007 0.011 Cu 0.035 0.07 0.015 0.13 S 0.006 0.007 0.004 0.018 Table 2: Heat treatment steel code Heat treatment GW PWHT: 610C-20h,f7c FF A: 865/880 C- 3h, WQ, T: 630/650 C- 5.5b, a/c SPWHT: 550 C-35h/615C-16h.ee JP A: 880C-8h,WQ,T: 650C-6h,a/c SPWHT: 620 C-26h,ffc HSST 915C-12h,a/c,A: 860C-12h,WQ, T: 635C-12h,l/c,SPWHT: 610C-40h,0e A: auslenitizc; WQ: water quench; T: temper, SPWHT: simulated post-weld heat treatment; a/c: air cool; f/c: furnace cool .64- Table 3: Transition temperature and upper shelf energy steel code GW JP FF HSST TT41J(°C) -73 -44 -57 6 USE (J) 183 196 200 125 Table 4: Static and dynamic yield strength at room temperature steel code GW JP FF HSST Cy (MPa) 517 457 532 474 o,tf(MPa) 666 549 637 578 Table 5: Values of initiation and propagation energy at upper shelf region steelcode GW JP FF HSST Ej(J) 60 60 62 50 Ep (J) 123 136 138 75 E, (J) 183 196 200 125 -65- .si*-*! «J5& v (c) Tig.x : ^restructure or tested steels, x260 W (b) JP (C} GH {d) HfiST -66. 240 200 -z'jpz'jr-: tri-z-- 160 .A' g. 120 • a VI 80 - —O- O-O pp - -i—I—I— JP A0 - •••Q-tl"Cl-" HSST —*—f—Jt— GW 1 "1 1 1 -r - • ! "1 -- T -150 -100 -50 0 50 100 150 200 250 3001 Temperature (C) Fig.2 : Impact Transition curves of tested steels -67- -89- L U ft 0 fKN) L 0 ft 0 CkN) ' — — I\J — — ro in CD Li w o «- «• M? •Art H- 0 O o H3 eg •0 o n o IT LC10 CW,') LOtlD CkN) OK U CO Ul C3 M| o, Ol CO u 11 1111111 I III I II I 11 otj I I 111 11 I I 111 I I I 11 111 X a o a as w u * rt co a en u »*rat a » ct M 800 - 4* » M' GH -+-+—I— JT 700 - •-n--Q-- HSST ! I" I 500 £00 -150 -100 -50 50 100 150 200 250 300 fenpernture (C) Fig.4 : Effect of test tenperature on dynamic yield strength •69- strengths at room temperature together with the corresponding static values. The variation of Ei and Ep with test temperature is shown in Figs.5 and 6, respectively. In the upper sliclf region, the energy consumed during the crack propagation process, Ep, is much higher than that required for the crack initiation process, E;. Values of Ej and Ep together with the total impact energy (E,) arc given in Table 5. E,- increased from 50 J for the reference steel (HSST) to 60 J for the advanced steels, At the same time Ep exhibited much greater increase, from 75 J for the HSST steel to about 130 J for the advanced steels. This indicates that the improvement in the upper shelf energy for the advanced steels over that of the HSST steel was mainly due to the increase in the propagation energy fraction. Another observation which deserves interest is the decrease in Ei with increasing testing temperature in the upper shelf range, Fig.5. 3.3. Fractography Fracture surface examination revealed that in the case of brittle fracture (lower shelf and transition regions) the fracture was by cleavage, while in the ductile fracture case (transition and upper shelf regions) the fracture was by microvoid coalescence. In the latter case, both large and small dimples were observed with no major differences between the four tested steels. On the other hand, different cleavage fracture features for the different steels were found, Fig.7. The cleavage morphology in the cases of the HSST and JP steels comprised facet-like arrangement with riverpattern s formed by cleavage lines and steps. The facet size can be related to the grain size in both cases. The SST steel showed continuous fracture plane, though changing orientation from grain to grain. However in the JP steel, the fracture plane experienced abrupt changes from one grain to the other. In the cases of the FF. and GW steels the cleavage fracture was characterized by smaller cleavage facets connected by tear ridges (quasiclcavage). The dimensions of the facets in the GW steel are comparable to the aciculai fertile grain size which constitutes the microstruct- ure of this steel. However, the facet size in the FF steel is probably related to a microstmct- ural unit finer than the prior austenite grain size. 4. DISCUSSION The results have shown that the advanced steels acquire better impact toughness than the reference steel (HSST steel) as expressed by their lower DBTT and higher USE values, Table 3. In addition, the advanced steels had higher strength in both static and dynamic conditions as seen above. In this way, these steels have good combination of strength and toughness compared to the HSST steel. This improvement in the toughness level can be thought of in light of the difference in chemical composition and microstructure between the four steels. -70- 100 -, 75 \t 50 5 25 1 r l t i - - 1 -150 •100 -50 0 50 100 150 200 250 300 Temperature (C) Fig.s : Effect of test temperature on initiation energy -71- 1G0 -, V" -150 MOO 0 50 iro' 1S^ 200 250 Temperature (C) 300 Fig. 6 : Effect of test temperature on propagation energy .72- Fig.7 : Scanning electron fractographs of tested steels, x380 (a) FF (b) JP (c) GH (d) U6ST -73- The advanced steels have relatively lower contents of carbon, sulphur, copper and phos phorus compared to the HSST steel. Carbon and sulphur have been known to have detrimen tal effects on toughness of steels. Increasing C content was found to raise the DBTT and decrease the USE [7]. Increasing sulphur content was found also to reduce the USE [8]. These effects can be explained as follows. C and S in steels form second phase particles (carbides and manganese sulphides). The density and/or size of these particles will increase with increasing C and S contents. Carbide particles participate in both cleavage and ductile dimple fracture. Crack initiation in cleavage fracture of steels occurs mostly at carbide part icles. The increase in density and/or size of carbides is expected to lead to more crack nuclei which act to decrease the cleavage fracture strength and hence raise the DBTT [9]. On the other hand when large and small second phase particles are present in steels, the ductile fracture process involves void nucleation and growth around large particles (MnS inclusions) and eventual coalescence via shear band localization from voids formed around the numerous smaller particles (carbides) [10], The increase in the density and/or size of the second phase particles, due to increasing S and C contents, is expected to increase the density and/or size of both the large and small voids. This will enhance the premature termination of the void growth process and, consequently, reduce the fracture energy as seen in Figs.2,5 and 6 above. The microstructure also is thought to have played a role in the toughness difference between the investigated steels. It was shown above that the different microstructures of these steels resulted in different features of cleavage fracture. While the HSST and the JP steels showed cleavage fracture appearance, the smaller facet size and the abrupt changes in the fracture path in the JP case are indications of a higher cleavage fracture strength. Such an improvement accompanies also the quasicleavage fracture features as those observed in the FF and GW cases. It is believed that microcracks in the case of the advanced steels were hindered by more effective barriers than in the case of the HSST steel. Thus the crack is forced to reinitiate repeatedly and more energy is expended in the fracture process; lowering the DBTT and increasing the fracture energy. As shown in the results, the difference in the impact energy between the advanced steels and the reference steel was mainly due to the difference in the fracture propagation energy. This is in agreement with previous investigations [10,11] which showed that the change in second-phase particle parameters (shape, volume, distribution) has a much larger effect on crack growth process as opposed to crack initiation process. This was explained in light of. the difference between the two processes [10]. Initiation is fundamentally a two-dimensional process. In such a case, crack blunting takes place effectively in the straight ahead direction so that only (hose parameters associated with this direction will be important. Propagation, however, occurs by linkage of voids which are spatially distributed, so that second phase particle parameters of all three dimensions must be taken into account [10], The result is that the change in these parameters will affect crack propagation process to a much higher degree than crack initiation process. The upper shelf of the Charpy impact energy curve is often viewed as a horizontal line -74- ,i.c., the USE docs not vary with the test temperature. The present results showed that the energy fraction representing the fracture initiation ,Ej, decreased with increasing test temper ature. This is in agreement with typical behaviour of the fracture toughness (Kic, Jjc) with test temperature. This could be understood since K|c and J|c are also concerned with fracture initiation process 5. CONCLUSIONS Instrumented impact testing was carried out for three advanced pressure vessel steels in comparison with a conventional pressure vessel steel (HSST). The main conclusions arc: 1. The advanced steels acquire much better impact properties (lower ductile-brittle transition temperature and higher upper shelf energy ) than the conventional steel. 2. The increase in the impact energy of the advanced steels is mainly due to the increase in the fracture propagation energy rather than the fracture initiation energy. 3. The improvement in the toughness of the advanced steels compared to that of the HSST steel arc related to the difference in chemical composition and fracture surface morphology. 4. The fracture initiation energy fraction of the impact energy decreases with increasing temperature in the upper shelf range which is in agreement with the behaviour of the fracture toughness Kic. REFERENCES [I] International Atomic Energy Agency," Analysis of the behaviour of advanced reactor pressure vessel steels under neutron irradiation", IAEA Tech.Rep.256 (1986). [2] Hawthorne,J.R., in "Treatise on materials science and technology vol.25", Ed Briant, C.L., and Banerji,SJC, Academic Press (1983)461-524. [3] Kawaguchi,S., ct al, Nucl.Engng Design 81 (1984)219-229. [4] Kussmaul,K., Nucl.Engng Intcr.(Dec.l984) 41-46. [5] ASTM E23-72, Book of ASTM Standards, Part 10 (1980). [6] Servcr.W.L., J.Eng.Matcr.Tcchol.,100(1978)183. [7] Gladman,T. and Pickering,F., in "Yield, flow and fracture of polycrystals", Ed.Baker, T.N., Applied Science Publishers, (1983)161. [8] Hcrtzberg,R.W.," Deformation and fracture mechanics of engineering materials",John Wiley & Sons (1989)365. [9] Tetclman,A.S., and McEvily.A.J. "Fracture of structural materials", Wiley(1967)268,523. [10] Howard,I.C. and Willoughby,A.A., in " Developments in fracture mcchanics-2", Ed. Chell,G.G., Applied Science Publishers (1983)39-99. [II] Ritchie, R.O. and Thompson,A.W.: MctTrans., vol. 16A (1985)233-248. -75- First Internationrl Spring School & Symposium on Advances in Material:. Science (SAMS 94) 15-20 March 1994 FRACTURE ANALYSIS OF SILICON NITRIDES LD. Abdelrazik, A.F. Bishay and A.M. El-Eslabi Nuclear Research Center, AEE, Cairo Egypt G.KUest and H. Nickel IRW, KFA, Jtilich, Germany. Abstract: Fracture analysis was performed on bend specimens of two grades of silicon nitride: Sintered silicon Nitride (S-S:3N4) and Hot Isostatically pressed Silicon Nitride (HIP-S13N4) The main objective of the fracture analysis was to find the fracture initiation point, to describe the nature of flaw leading to fracture and to compare between the measured bending strength and local fracture strength derived from the fracture mechanical evaluation. The stress intensity factor (Kc) was calculated for the flaws supposed to be the fracture origins and compared with the fracture toughness evaluated using the Vickers indentation technique. Examination of the fracture surfaces of S-Si3N4 revealed three types of flaws. These were pores,porous regions and inclusions. The most frequent flaw type was the porous region. The porous regions were originated in the green body microstructure as density inhomogenity and may have evolved during sintering into its final state. Examination of the fracture surfaces of HIP-Si3N4 specimens revealed the presence of four types of flaws. These were surface flaws, edge flaws, porous regions and inclusions. The most frequent flaw type was the surface flaws. The surface flaw were originated during surface finishing of specimens. The stress intensity factor calcualted for the flaws, supposed to be the fracture origins, increased with the flaw size to certain limit for both S-SioN^j and HIP-SioN^. The measured fracture strength showed poor agreeent with the fracture strength calculated from Griffith equation. 1. Introduction: Silicon Nitride are among the most important materials developed for high temperature structural and wear resistant applications, because of the good combination of mechanical, thermal and thermo- mechanical properties. These properties include high strength at high temperatures, good thermal stress resistance, good oxidation resistance and corrosion resistance [1]. At lower temperatures, the strength behavior of structural ceramics is characterized by brittle failure i.e. tendency to fail catastrophically by growth of single crack that originates from a very small flaw [2J. These flaws, inherent in the microstructure, interact with the stress field to increase the local effective stress. The result of stress intensification is that the flaw rapidly propagating across the -77-- specimen causing fracture [3]. The stress intensity is directly related to the size and shape of these flaws, with the result that fracture strength depends inversely on flaw size. Thus to increase the reliability of ceramic parts; one must (1) control the flaw population in terms of size, shape and distribution, (2) increase the toughness (resistance to crack propagation) of the material. The aim of this work are to, find the fracture initiation point, describe the nature of flaw leading to fracture, and to compare between the measured bending strength and local fracture strength derived from the fracture mechanical evaluation. 2. Experimental Procedures: 2.1. Materials: Sixty broken specimens of HIP-SigN^ [ASEA, Sweden] and sixty broken specimens of S-SigN^LGTG-USA] were supplied by Hoechst Ceram Tec Company. The samples had the nominal dimensions of 4.5 x 3.5 x 50 mm. The samples were fractured by four-point bending with cro? i head speed of 0.5 mm/min. in air by Hoechest Ceram Tec Ccaupany. The bend fixture used had an inner span of 20 mm and an outer span of 40 mm. 2.2. Proctography: Only 15 specimens were selected for fractographic study from each grade. For each grade, the specimens were ranked by strength and then divided into three groups consisting of the five highest strength specimens, the five lowest strength specimens, and five specimens taken from the middle strnegth distribution. The selected specimens for fractography were examined to identify the fracture origin and the type of flaw leading to fracture. The fracture surfaces associated with each specimen were examined using a Lietz stereo-light microscopy with maximum magnification of 24x. All the fracture mirrors were identified. The fracture surfaces were then examined using SEM with a magnification of lOOx, lOOOx and 2000x to identify the fracture origin as well as the type flaw present. 2-3 Fracture Toughness Measurement: The fracture toughness of the two grades of SigN^ was measured using Vickers indentations technqiue. The sample segments of each grade were mounted and then polished by using diamond pastes 15,6,3 and 1 urn respectively to produce optical finish and stress free surface. Vickers indentations were made by using a standard Vickers indenter with apex angle 0= 136*. AJl indentations w;re made in air at ambient temperature with a steady loading rate that required 10 sec to reach full load and 10 sec at load. The indents were made at loads ranging from 30 to 392.4 N. The -78- number of indents at each load made in a specimen ranged from 5 to 10 indents. The indentation half diagonal lengths and radial crack lengths were measured directly using a Lietz optical microscope with caliberated eyepeice graticule and a magnification appropriate to the size of the indents (typical X 200 to X 1000). The average radial crack lengths, and the half diagonal lengths were used in the indentation fracture toughness equations. 3. Results and Discussions: 3.1 Specification of fracture origins: At room temperature, 60 specimens from each grade were tested in four-point bending to determine the fast fracture strength. The bending tests and the statistical variation in fracture strength were performed by Hoechest Coram Tec company. The fracture strength of S- SioN^ varied from 375 MPa to 789 MPa with an average strength of 651 MPa, Weibull modules of 11.9 and a standard deviation of 73 MPa. While, the fracture strength of HIP-Si3N4 varied from 477 to 775 MPa with an average strength of 648 MPa, Weibull modulus of 14.2 and a standard deviation of 52 MPa. Examination of the fracture surfaces of S-SigN^ specimens revealed the presence of three types of flaws. These are pores, porous regions and inclusions. As shown in Table 1, the most frequent flaw type, was the porous regions. The typical dimensions of the pore which led to failure were of the order of 50 um or more. Typical example of surface and subsurface poras or pores region are shown in Fig. 1 and 2 respectively. In most cases rod-like and prismatic like grains of P-SigN^ were developed in the hole (See Pig. 3). Fracture has initiated from a large surface pore or porous region in the lowest fracture strength specimen (375 MPa). The maximum extension of porous region was about 400 um. Failure initiating porous regions varied in size, ranging from 21-275 um long and 20-395 um wide, circular, semi-circular and ellipsoidal in shape. The presence of such large porous regions suggests that full densification of the bulk material did not occur during sintering. The results of this study is confirmed by the findings of Buljan et al [5] and Pasto et al [3] for S- SigNj containing 6 wt% Y2O0 and 2 wt% A^Og as sintering additives and Govila [4] for S-SigN^ containing various oxide additives (Ce O2, Mgo, Zr 02 and SrO). Buljan et al [5] and Pasto et al [3] studied the origin of the porous regions. They found that this flaws originated in the green compact body microstructure as a density inhomogenity. This density inhomogenity is likely caused by local differences in the density among the agglomerates, which typically occurs in dry processed - SigN4 and which are not completely eliminated during pressing. -79- Tabic (I): FrMO£raphy summary Table lor S-Si.N,. Samp c Defect Defect Dimensions DcfecuofKxHnaies Measured (Frame) urn (center) / mm fracture No. type Y Y AZ X Y z strength.MPa 1 surface 275 0 395 20.49 n.ixi 2.98 375 porous region 2 Surface pore 152 0 81 16.27 t).(X) 2.25 551 3 Surface pore 84 0 120 31.28 0.0(1 1.16 555 4 Surface porous 36 0 60 30.06 0.00 3.17 571 Region-nearby 79 0 118 5 Subsurface pore 104 16 70 11.86 0.06 0.33 380 6 I Subsurface pore 73 20 55 19.60 0.054 1.88 630 7 Porous region 35 0 50 19.52. 0.00 fl.64 650 on the surface 8 Subsurface 106 22 86 31.47 0.07 1.40 626 Mrous region V Subsurface 46 10 54 3I.X9 0.03 0.52 646 TOK 10 1 Subsurface pore 82 38 44 16.53 0.06 0.33 635 II !Subsurfac e 84 37 35 18.46 0.06 0.92 703 i nclusions 12 S ubsurface 24 3 29 23.34 0.01 0.53 727 iiiclusio n II I S 13 ii 47 3(1 20 22.03 0.05 0.20 716 IV 14 S urficc pore 45 0 45 19.37 0.00 0.58 750 •iih inclusion 35 0 30 15 S urface pore 25 0 28 18.75 0.00 3.99 789 - -80. 08U-BJ* 8«<99 1RW Figure 1 , Surface pore observed at fracture origin of S-Si,N, bend specimen I8KU x>84er. leu ar9 e«49a IRU Figure 2, Subsurface spherical pore observed in fracture surface ofS-Si,N4. Figure 3 , SEM micrograph of porous region observed at fracture origin of S-Si,N, specimens showing that the porous region is filled with both rod-like grains and some prism like grains. -81- Failure was initiated in three samples out of fitecn from metallic inclusions. Energy dispersive X-ray analysis of these inclusions usually show high Si plus any several element species, principally, 0, Na, Al, S, CI and Ca suggesting their pickup during powder preparation, processing and its handling to flnishcd shape [3]. A typical example of subsurface inclusions is shown in Fig. 4 and its ED AX analysis is shown in Fig. 5. Examination of the fracture surfaces of HIP-SioN^ specimens revealed the presence of four types of flaws. These are edge flaw, surface flaws, porous region and inclusions. As shown in Table 2, the most frequent flaw type was the surface flaw. Most critical defects were located in the vicinity of sample surface. The absence of large pores or porous regions in the HIP-SigN^ was probably a consequence of the extcrmcly high pressure in the processing of this material. Although the edges of the samples were chamfered, three samples were failed from edge flaws. The fracture occurred from edge flaw in the lowest fracture strength speimen. Failure of samples due to group of inclusions took place in three cases out of fifteen. Energy dispersive X- ray analysis of these group of inclusions showed several element species, principally O, Na, Al, CI and K. A typical example of subsurface group of inclusions and their EDAX analysis are shown in Figs. 6 and 7 respectively. The surface flaws seems the most difficult one to be detected and prcciesely specified. The surface flaw included machining damage, penetration flaw and stepped flaw. In one sample only, penetration flaw was found at the fracture origin of HIP-Si^N^ . Penetration flaw consists of surface damage in a plane that is being intersected by the fracture surface. As, a result, this flaw appears in the fracture surface as a linear defect extending from the surface toward the axis of the specimen [6], Penetration flaw may be created during surface finishing of the specimen. They are believed to be the result of deep scraches oriented at an angle 90* to the fracture surface [6]. Example of penetration flaw is shown in Fig. 8. A machining defect was found at the fracture origin of HIP-S13N4. Failure in three samples was initiated from stepped flaws. Stepped flaws were found at or near the surface and have a step like appearance. This flaw has occurred during surface finishing [3]. A typical example of machining damage and stepped flaw are shown in Fig. 9 and 10 respectively. Stepped flaws and penetration flaws were also found at. the fracture origins of HIP AI2O3, HP-SiC and HP-Si3N4 16]. 3.2 Fracture toughness measurement: The fracture toughness of S-Si^N^ and HIP-Si-jN^ was measured using the identation technique. The validity of the use of the Vickers Table (3): Fractography summary Table for HIP-Si.N, .S-imj le Delect Defect Dimension* Defect coiTodinaics Measured (Frame) urn • (center)/mm fracture No type Y„ Y„. a. X Y z strcngth.MPa 1 SurfK-e 40 II 65 34.60 o.oo 1.00 518 Flaw 3 Edge failure • • 36.00 0.03 4.55 477 3 Briud agglome< 411 (1 511(1 35.10 o.m 4.2V 545 miiwofinclusi ons near surface 4 I Edge Flaw due 65 (1 71) 33.6(1 0.00 4.29 572 1 u> inclusion ... 5 .Edgul'aituni •- •• 17.80 0.00 4.26 582 (weak region) fi Mxhining 55 0 1(10 34.80 0.00 1.70 605 damage 7 Surface flaw 55 0 100 34.0 0.00 1.70 651 )) Subsurface 68 >2 80 14.40 0.04 1.30 602 roraus region «> Surface flaw 25 (1 85 25.30 O.(X) 4.37 653 10 Stepped flow 61 II 245 35.50 (UK) 1.55 633 II Stepped flow 55 0 237 18.76 0.00 1.60 721 12 Surface flaw 25 0 98 18.75 0.00 3.93 1 722 1.1 1Vnenirau'o n flaw 65 0 65 32.10 0.00 0.70 723 14 J ubsurface group 96 58 35 30.90 0.08 4.2 7011 0 f inclusions 15 S teppedflaw 32 0 55 19.50 0.00 0.60 775 -83- Figure* , Subsurface inclusion observed at fracture origins. CURSOR Figure 5 , Energy dispersive X-ray analysis of the inclusion observed at fracture origin of S-Si,N« . -84- Figure 6 * Group of inclusion observd at fracture origin of HIP-Si3N«. QCI N A SYl S C K C A L I I L A. CURSOR Figure 7, EDAX analysis of inclusions observed at fracture origin ofHIP-Si,N,. -85- Figure 8, SEM micrograph of penetration flaw observed at fracture origin of HlP-Si^. Figure 9, SEM micrograph of machining type defect (machining damage) observed at fracture origins of HIP-Si,N4. Figure lb, SEM micrographs of stepped flaw observed.at fracture origin of HlP-Si3N4. -86- indentation technique to measure the fracture toughn.ess Kc for both materials was checked by evaluating the dependence of measured surface crack length, C, on the applied load, P for loads ranging from 29.4 to 294.3 N. This load range was chosen to ensure that C > 2a (a is impression half diagonal) without chipping. This ensures a clearly visible well developed cracks influenced only by the far field stresses [7]. The dependence of the indentation crack length, C, on the test load, P, is shown in Fig. 11. Theory requires that P is proportional to C^'2 [7]. A slopes of 1.54 and 1.492 were found for both S-S13N4 and HIP-S13N4 respectively. These slopes are very close to the theoretical value of 3/2. Therefore Kc can be calcualted for both materials using Anstis Equation [7] and Evans and Charles Equation [8]. The results of the S-Si3N4 and HIP- S13N4 using these two equations are shwon in table 3. Table (3): Mean Vickers indentation fracture toughness for SWHgN4 and mp-si3N4. Equation used S-Si3N4 HIP-Si3N4 MPa. m 1/2 MPa. m 1/2 Anstis et al Eq. 5.7 ±0.1 4.2 ±0.1 Evans and Charles Eq. 6.73 ±0.2 5.00 ± 0.1 2L3 Calculated fracture toughness: The stress intensity factor, Kc values were calculated for flaws supposed to be the fracture origins through fractographic observation from the formula. Kc=YopVc^ (1) Where Y is the geometrical correction factor (1.28 [10] for semi- el lipitical surface flaw) &p fracture strength and Cq is an equivalent crack length. The equivalent crack length is given by yTb, where r and b are semiminor and semimajor axes of flaw. The Kc values were calculated for specimen with surface and well defined pores or porous region for S-SioN4 and well defined surface flaw for HIP-Si Kc seems to increase with increasing crack length. This agrees with the findings of Ferber [9] for both S-Si3N4 and HIP-SigN4 , and -87- 1.7 1.9 2.1 2.3 2.5 2.7 logc/Jm Figure 14., Plot of indentation load and half radial crack length for indented specimens of HIP-Si.,N4. -88- El-Aslabi [11] for HP-SigN^. The increasing of fracure toughness with crack length may be attributed to two reasons: (i) The flaw geometry which is not taken into consideration in these calculations, (ii) These two materials may experience R-curve behavior. The average value of the calcualted fracture toughness of S- Si3N4 was 5.73 MPa. This value was comparable to the measured fracture toughness value derived from Vickers indentation technqiue using Anstis et al. Equation. The average value of calcualted fracture toughness of HIP-Si3N4 was 6.04 MPa. This value was higher than measured Kc values using indentation technqiue. This method of calculation seems to over-estimate the fracture toughness of HIP- Si3N4. This could be in part attributed to the uncertainty associated with the flaw size measurements. Another possibility is that a part of the flaw may act as critical defect when the fracture originating at penetration flaw, group of inclusions or stepped flaw. Kirchner et al [6] found that the theoretical flaw size predictions for fractures originating at stepped flaws and at penetration type flaw, were much smaller than the experimentaly determined ones. 3.4. The flaw size strength relationship: The flaw size-strength relationships for S-Si3N4 and HIP-SigN4 were examined by plotting the fracture strength, Of as a function of flaw size, d, (crack depth) for surface and subsurface flaws. In the case of subsurface flaws, the corresponding values of flaw depth were used to correct the fracture strength values. The results are represented in Figs. 13, 14 for both materials, along with the predicted behavior (solid line) determined from Griffith equation. y2 Of- (l/Y) Kc d- (2) ue The assumed values for Y and Kc were 1.28 [11] and 5.7 MPa m " 2 respectively for S-Si3N4, and 1.28 and 5.00 MPa.m^ respectively for HIP-Si3N4,. As shown in Figs. 14 and 15, the data exhibited considerable scatter and did not correlate well with the predicted Griffith relationship. In particular, there was a tendency for the measured fracture strengths to lie above the predicted values when the defect exceeded 100 urn for S-Si3N4> and 40 urn for HIP-Si3N4, A similar result for S-Si3N4, and HIP-Si3N4, was reported by Ferber [9] and Hoshide et al [10] for S-Si3N4, . These results support the idea that S-SigN4, and HIP-Si3N4, may experience R-curve behavior. -89- Kgurei2, Relation between equivelant crack length, Ceq. and critical stress intensity factor 100 -eq. »/»v» m Figure 13, Relation between equivalent crack length, Ceq. and calculated stress intensity factor. -90- -i 1 r « Surface Flaw a Subsurface Flaw 2.1 —-Griffith Eq. _i I I i L U M l.« *4 2-0 2.2 log d. JJm Figure 14, The relationship between flaw size and fracture strength forS-Si,N«. i3 0 Surface Flaw o Subsurface Flaw -—Griffith Eq- 13 1.4 1.8 '•• 20 2.2 logd,/"" Figure 15, The relationship between fracture strength and equivalant crack length for HIP-Si3N4. -91- 4- Conclusions: (i) Strength of S-SigN^ is controlled by microstructure defects, mainly porous regions. The porous region may originate in the green body microstructure as density homogeneity and may evolve during sintering into its final state, (ii) Strength and uniformity of S-Si3N,j can be improved through microstructure control (porous region) and further improvement are possible through control of other flaw types such as metallic inclusions, (iii) Strength of HIP-Si^N^ is controlled by surface finishing. (iv) Strength of HIP-Si3Nj can be improved by improving surface finishing, (v) The measured fracture strengths for both materiala did not correlate wellwith the strength calcualted for the flaws supposed to be the fracture origin using Griffith Equation. References (1) G Ziegler, J.Heinrich and J. Wotting. J. Mater. Sci., 22 (1987) 3041. (2) D.B. Marshall and J.E. Ritter: Ceram Bull., 66[2] (1987) 309. (3) A.E. Pasto, J.E. Neil and C.L. Quekenbush: in "Ultra-structure processing of ceramics, Glasses and Composites" edited by D. Larry and R. Hench, John Wiley & Sons, Inc. New York (1984) P. 478. (4) R.K. Govila: Int J. High Technol. Ceram., 3 (1987) 179. (5) S.T. Buljan, J.T. Neil, A.E. Pasto, J.P. Smith and Zilberstein: in "Non oxide Technical and Engineering Ceramics". Edited by S. Hamphire, Elsevier Press, London (1986), P. 409. (6) H.P. Krichner, R.M. Graver and W.A. Satter: Mater. Sci. Eng., 22 (1976)147. (7) C.R;,Anstis, P. Chantikul, B.R. Lawn and D.B. Marshall: J. Am. Ceram. Soc, 64 [9] (1981) 533. (8) A.G. Evans and Charles: Ibid, 59[78] (J976) 371. (9) M.T. Ferben Unpublished Work. (10) T.Hoshide, H. Furuya; Y. Nagase and T.Y. Amoda. Int. J. Frac. 26(1980)229. (11) A.M. El-Eslabi: Ph.D. Thesis, Alexandria Univ. (199CV -92- First International Spring School & Symposium on Advances in Materials Science (SAMS 94) 15-20 March 1994 FORMATION OF ACSTENITIC STAINLESS STEEL BY SOLID STATE REACTION M. SHERIF EL-ESKANDARANY1 AND H.A. AHMED2, I. Mining and Petroleum Engineering Department. Faculty of Engineering, AI-A/.har University, Nasr City, Cairo, Egypt. 2. Mining. Petroleum ami Metallurgical Engineering Department, • Faculty of Engineering, Cairo University, Giza, -gypl ABSTRACT A new process for synthesis of austenitic stainless steel powder of type Cr to Ni as 18:8 nt room temperature is reported. This process depends on a solid state diffusion between the elemental powder nf Fe, Cr and Ni using rod milling technique. The alloy powders have been characterized as a function of the diffusion time (milling time) by means of X-ray diffraction, optical microscope, scanning and transmission electron microscopes and differential thermal analysis. The results have shown that a homogeneous amorphous phase of F^Cr.^Nig alloy powders has been formed after 1080 ks (300 h) os' the milling time. The formed amorphous phase consists of uniform (spherical-like morphoi- gy) ultra fine (almost 1 |im in diameter) powder particles. The internal structure of the end-product of the alloy powder shows a very fine strucih t: free from any defects or strains. The thermal stability indexed by the crystalnzaucn temperature, Tx and the r i: r ciillulpy change of cryslallizaliou. AH of ' 74*- |XNi alloy powders show that the fabricated austcnitic stainless steel amorphous alloy is very stable at the high temperature, as high as 1240 K. Above this temperature (K) an amorphous-to-crystalline phase transformation takes place so that all the amorphous phase of the austenitic stainless steel has changed to fec-structure with AH - 16.88 kJ/mol. Key Words: Solid state reaction, Diffusion, Mechanical alloying, Rod milling, Lamellar- structure, Morphology, Defects, Solid solution, Amorphous phase, Austenitic stainless steel, Powder metallurgy, Structure Morphology, Crystallization - 93 - 1. INTRODUCTION The first example of two pure crystalline metals reacting to form a single-phase amorphous alloy was given by Schwarz and Johnson 111. In this experiment, thin films of pure gold and lanthanum, a few tenths of a nm in thickness, were fully reacted at 343 K within a few ks. Two requirements have been proposed for the solid-state amorphizing reaction: (a) the two reacting metals must have a large negative heat of mixing and (b) the two metals must vastly different diffusivities in each other and in the amorphous alloy to be formed. The first condition ensures that a thermodynamic driving force for the reaction exists. The second condition ensures that the amorphous alloy will form in preference to crystalline intcrmetallics, which have lower free energies. This kinetic selection of the reaction path is possible because one species diffusing in the other and in the amorphous alloy is sufficient for the solid state amorphizing reaction. The micro-mechanism of the solid-state amorphizing reaction in multilayers was studied by Rutherford back-scattering spectroscopic (RBS) marker experiments on a Ni/Zr diffusion couple [2] and transmission electron microscopy on Co/Zr [3]. These experiments showed that the smaller atom in the diffusion couple is the moving species. The growth rate of the amorphous layer, which is formed at the interface of the two atomic species, turned out to '•- determined by the diffusion coefficient through this amorphous layer. A different attractive technique for preparation of amorphous materials at temperatures below the crystallization temperature is mechanical alloying (MA) of elemental crystalline powders or the grinding of the crystalline alloy, that is to say, mechanical grinding (MG) [4J. White [5J first reported the .raorphization of Nb Sn by MA and then Koch et al. |6J established the synthesis of amorphous Ni No powders by milling a mixture of pure elements. In the typical cxperi.i.rat for preparation of an amorphous Ni Nb alloy, elemental powders of Ni and Nb were milled for about 40 ks in He gas using a Spex Mixer/Mill Model 8000 and steel balls as milling media. The end-product of amorphous Ni Nb alloy powders was contaminated with oxygen (0.43 wt.%) and Fc (4.11 wt.%). The MA process is based on the solid-state amorphizing reaction. It is worth noting that the composition range for which amorphous metallic alloys produced by MA is larger than that for melt quenching, where the glass-forming concentrations are centered around the cutectic composition. One significant potential technological attraction of MA process is that the resulting amorphous powders may subsequently be consolidated by hot pressing at temperatures below the crystallization temperature of the amorphous alloy, thereby offering the prospect of manufacture of bulk objects made -94- from amorphous metallic allocs, wi'hout loss of their uni(|iie physical properties, a circumstance which is impossible using mcll-qticuching techniques |7|. In fact, the MA process dates back to 1970 when Benjamin |S| employed the ball- milling technique (UM) for producing homogeneous cont|K>siic particles with intimately dispersed, uniform inicrnal structure. So far, the MA process has been used for preparing many dispersion-strengthened alloy powders |9-)5|. Comparable to MA process, amorphous alloys can lie prepared by MG. Ycnnakov et al. H| first reported that the mechanical attrition of crystalline intcrmctnllic comfiounds of Y and Co led to the formation of an amorphous alloy powder. The MA process synthesizes amorphous alloys by reacting elemental crystalline powders, being accompanied with negative heat of formation. In MG process, however, crystalline alloy orcom|)ound powders are transtormed into amorphous solid powders by destroying the periodical long-range order of atomic arrangement without compositional changes [ I6|. MA and MG processes are the reactions going to thcniiodynamically opposite each other The MA process via UM technique has been considered as a practical solution for amorphizing the systems which are difficult or impossible to be obtained by conventional melting and/or casting techniques, e.g. Al-Ta 117| and Al-Nb 118| binary systems. More recently, fcl-Eskandarany ct al.| 19, 201 have -cportcd the first novel technique for the formation of metal nitride by milling elemental TM (TM; Ti and Fc) powders under flow of nitrogen -it room temperature using the reactive ball milling (KBM) process. The KBM leads to the formation of high stable amorphous AITa contains 15 at.% of nitrogen |21, 22| and has been also used for formation of superconducting Nb-N alloy powder |23|. In l«W0, bJ-Hsk;mdarany ct al. |24| reported another novel technique for producing amorphous Al Ta_ alloy powders front pure aluminium and tantalum powders by a method called rod-milling. This unique technique has been accented as a successful method for producing amorphous alloy powders with low degree of contamination, that is to say. iron contamination and high thermal stability through solid-state amorphi/.ing reaction |25|. The application of this new technique has led to the formation of many amorphous alloy powders 126-351. In the present work amorphous aiistcnilic stainless steel Itc-Cr.Ni) lias been prepared by rod-milling technique via MA route ;il room lcni|>craturv. This study has - 95 - been addressed in part to offer an industrial application of the MA process for producing an alloy which hns great interest in industry, material science and engineering points of view. The morphological, structural and (hernial stability behaviors of amorphous Fe74Cr1HNiH alloy powders are presented. Moreover, the mechanism indexed by the mode of amorphization and crystallization for Fc_.CrluNiualloy powders will be discussed. As far as the authors know, this is the first report for preparing austcnitic stainless steel by solid state reaction using rod-milling tcchniqu... 2. EXPERIMENTAL 2.J. Materials In the present work, a high purity (99.99 %) elemental powders of Fc, Cr and Ni metals have been used for the mechanical alloying (MA) experiments. The average particle size of Fc and. Cr and Ni elements was 40 and 80 jun, respectively. The metal powdrrs were scaled under high purified argon gas (O, and H,0 were less than 10 ppm). 2.2. Type of Mill The rod mill that used in this s\udy has been designed so that its length (250 mm) is greater than its diameter (120 mm) and the rods (SUS 304, 10 mm in diameter) have been cut to lengths (200 mm) less than the full length of the shell. The movement of the rods (10 rods) inside the shell was directly observed through a Ihicl. and transparent plastic plate sealing the window of the shell. This observation has shown that the milling process occurs by the line contact of rod-powdcrs-rod extending over the full length of the shell. More details of this mill can be found in a previous paper [24], In the present work, the rod milling were carried out by mounting the shells on a rotator running at a rate of 1.4 s . 2.3. Sample Preparation The starting materials of I!T MA process were placed in a glove bag filled with Ar gas, together with the milling tools. The elemental powders of Fc. Cr and Ni metals were blended using a ceramic, mortar and pestle in the &o\c bag lo give the desired composition of Fc74Cr(J Powder X-ray diffraction (XRD) data were obtained with Philips-Norelco operating in the 0-28 scan modes for CuKa radiation at 35 kV and 20 mA. The sample was placed on doubly sticky transparent upc mounted onto a microscope slide and placed in the area irradiated by X-ray. 2.4.2. Thermal Analyses The crystallization behavior of amorphous Fe_.Cr.gNL alloy has been monitored by differential scanning calorimetry (DSC) at a heating rate of 0.33 Ks" in purified argon gas. In the DSC experiments, an amount (about 20 mg) of milled powders was inserted in a sample platinum pan. The lining of the pan was coated with ct-AI.O. to avoid any reactions between the samples and the platinum pan which could give false exothermic peak(s) especially at elevated temperatures. Further, a certain amount of a- AI.O, was put into the reference platinum pan. In the all measurements, the samples were heated up to 1300 K and cooled down about 500 K. Then the samples were reheated as second heating run to find the base line. X-ray diffraction was used to confirm the exothermic peak. 2.43. Metallography The metallography of the alloyed powders has been studied by a scanning electron microscope (SEM) equipped with an electron probe microanalyzer (EPMA) operated at 20 kV and light optical microscope. In these experiments, a smaii amount of the alloyed powders was placed in a plastic holder (epoxy) and wot ground carefully using fine silicon carbide emery papers (S0G-I500 mesh). The grinding process was performed using paraffin as a lubricant to prevent particles of emery being embedded in the surface of the specimen. Then, the samples were polished on lapping tape (2000-6000 mesh) wetted by kerosene. Finally, the specimen was polished on silk cloth using a solution of white wax in benzene. Care was taken until a satisfactory surface was obtained without any distortions. Then, the specimen was etched using HC1 (0.1 N) activated with five dropsofH,02(10%). 2.4.4. Morphology -97- The morphology of the alloyed powders were examined by a 20 kV SEM and a 200 kV transmission electron microscope (TEM). For the SEM observations, a small amount of the powders was put on a copper holder coated by carbon dutite. Whereas the samples for TEM observations were prepared by crushing the powders under cthanol. The samples had been dried before they were mounted on a copper microgrid. 2.4.5. Chemical analyses The chemical analyses were performed using ICP interfaced to a microcomputer in order lo detect exactly the contents of Fe, Cr and Ni. For these analyses, an amount of the milled powders (about 30 mg) was dissolved in a leaching solution [HF(1N), 2 ml+HnS04 (6N) 10 ml+HNC<3 (6N) 40 ml+H,0, (36 wt.%) 2 ml] using a platinum container (100 ml in volume). Then, the solution was heated slowly until the powders dissolved completely without any precipitations. The solution was diluted with pure water in a volumetric mass flask (100 ml in volume). The ICP was calibrated by Fe, Cr and Ni standard solutions before analyzing the sample. Each component for each sample was analyzed five times for more accuracy. 3. RESULTS 3.1. Structural change with the milling time X-ray analyses were performed in order to understand the total stvueture of Fe_.CrlcNL alloy powders at ihe several stages of the rod-milling process. Figure 1 displays the XRO patterns of Fe Cr Niu alloy powder after selected MA times. After 2 ks of the MA time (initial stage) the mill discharge consists of bcc-Fe, bcc-Cr and fec- Ni in the elemental form. The Bragg-pcaks of pure Fe, Cr and Ni decrease in their intensities and become board at the end of the initial stage of MA time (86 ks). After 173 ks of the MA time, all the Bragg-peaks for bcc-Cr and fcc-Ni crystals have completely disappeared. In addition, the Bragg-pcaks for bcc-Fc crystal have become broader and their positions have shifted remarkably to the low angle's side, indicating the formation of bcc-FeCrNi solid solution. It is worth notisi,; that after 360 ks of the MA time (intermediate stage) these peaks become broader with diffuse haloes, suggesting the formation of an amorphous phase coexisting with bcc-FeCrNi solid solution, as shown in Fig. 1. At the final stage of the MA time (1080 ks) the bcc-FeCrNi solid solution phase has been completely transformed to a homogeneous amorphous phase, characterized by diffuse haloes and smooth peaks. We should emphasis that this amorphous phase has not been changed to any other phase(s) even for the samples milled for longer time as long as 1800 ks, as presented in Fig. 1. -98- i i i i i i i i i i I i i' •1 I I I I I I I 1 I I II rod-milled Fe^Cr^Ni, fcc-Ni (220) alloy powders = — o 12 ks C3 J3 1 80 70 60 50 -*— 2 0 (degree) Fig. 1 Effect of the MA time on structure of rod-milled Fe Cr Ni8 alloy powders. 0.2880 i i i 11111 i i i 111 ii| i i i r i iiy Amorphization starts § 0.2878 <_, 0.2876 C 5 0.2874 / C O 0.2872 U 8 0.2870 '•£3 tS 0.2868 a -Fe powder 0.2866 ' ••' • ••••»•' IC 100 1000 Milling time (ks) Fig. 2 Effect of the MA time on the lattice constant, a , of bcc-Fc during rod milling of elemental bcc-Fe. bcc-Cr and fcc-Ni powders. -99- The lattice constant a„, of rod-milled Fc_ ,CrloNi. alloy powders estimated from U ll\ lo o XRD measurements arc shown in Fig.2 as a function of the MA times. Obviously, the bcc-FcCrNi solid solution expands with increasing the MA times, characterized by a monotonically increase in a to get a value of 0.2878 nm after 360 ks of the MA time. We should emphasis that this valus is larger than that of pure Fe crystal (0.28664 nm), suggesting an interstitial solubility of Cr and Ni in Fc. 3.2. Morphology and metallography changes with the milling time There can be no more fundamental characteristics of a powder than the size and the shape of the individual particles. SEM technique was used to follow the change in the morphology of mechanically alloyed Fc_ Cr Ni alloy powders at different MA times. Detailed SEM observations after several stages of the MA process are presented in Fig3. At the initial stage milling (2 ks) the powder particles that contain a mixture of elemental Fe. Cr and Ni have random size with block-like morphology, as shown in FigJ(a). After few ks of the MA time (22 ks) the starting elemental powders of the three elemental metallic powders have agglomerated to form composite particles of about 600 urn in diameter with plate-like morphology, as shown in Fig.3.(b). At this stage of milling, the powder particles have thick layered-structure morphology with random distribution, as presented tn Fig.-4(a). Increasing, the MA time (43 ks) enhances the shear and the impact forces generated from the rods, leading to the formation of composite powder particles contain thin or narrow layers of the elemental powders in an excellent arrangement and the individual layer has width of about 10 u.m or less, as displayed in Fig. 4(b). During the progress of the MA process (173 ks), the assemblage particles have continuously disintegrated to form powder particles that are almost regular in shape (globe or spherical-like morphology) and size( 10 to 20 urn), as shown in Fig3(c). The metallographic examination of the cross section for polished and etched particles after this stage of milling shows that the layered-structure morphology has already disappeared, and the individual particles have not any details in their polished surfaces, as shown in Fig. 4(d). At the end of the MA process (1080 ks) the powder particles are fairly uniform in their sizes that are vary from 0.5 to 1 pm in diameter, as shown in Fig.3(d). Figures 5 and 6 summary the SEM observations of mechanically alloyed Fe Cr Ni alloy powders at t'e several stages of milling . Obviously, the solid state reaction performed by the MA process can be classified into three stages, that is to say, -100. J g g s. u s. 00 U •o "8 Q. G 00 u2 c/5 CO -101- Fig. 4 Cross-sectional view of rod-milled Fe, Cr,JN^ alloy powder particles after (a) 11 ks. (b) 43 ks and (c) 173 ks of the MA time initial, intermediate and final stages. In the early stage the elemental powders of Fe, Cr and Ni have agglomerated and grown in size as a result of the repeated cold welding. During this stage of milling, the powders vary widely in size from SO to 800 (tin, as presented in Fig. 5. During this stage, the thickness of the layers in the formed composite powder particles increases dramatically and widely distributed, as showo in Fig. 6. At the subsequent intermediate stage (second stage of milling), the agglomerated powder particles are subjected lo a continuous disintegration with fragmentation to form finer powders with size less than 100 |im in diameter. Furthermore, this stage of milling provides powder particles with narrow size distribution, as shown in Fig. 5 In parallel, the thickness of the metallic layers decreases at this stage, becoming uniform with narrow thickness distribution, as displayed in Fig. 6. Towards the end of the MA process (the final stage) all the powders are uniform and homogeneous in size (Fig 5) without any structure details (Fig. 6).' -102- 800 | 1 I I I llll| LI II IIM| 1—I I I llll| ' I fi 600 >«n• « N CA 400 O J V 0 ' ' ' ' ' I Milt 0 # I 1 10 100 1000 Milling time (ks) Fig. 5 Particle size distribution of rod-milted Fe Cr Ni alloy powders as a function 7474 " 1lu8 t8t of the MA time. 40 ; I I I I llll| I III Hll| I I I I llll| I I I I 30 s % 20 lj £ io §- P . 0 0 ill null •—i i i mil ^ -j| j| 0 1 10 100 1000 Milling time (ks) Fig. 6 Effect of the MA time on the layer thickness of rod-milled Fe_,Cr„>NL alloy 74 18 8 ' powder particles. -103- Detailed TEM analyses were performed in order to observe the structure change of the rod-milled Fe_.Cr Ni„ powder during the above mentioned stages of milling. The bright-field images (BFl) and/or the dark-field linages (DR), and the selected-area diffraction patterns (SADP) of mechanically alloyed Fe_ Cr Ni powders after selected MA times are shown Rg.7. Figure 7(a) shows the BR and corresponding SADP of powders taken after 22 ks of the MA time. The powders are a mixture of polycrystalline of Fe, Cr and Ni that have grain boundary fringes and dislocations in the boundary. The SADP taking at the center of this micrograph showing a sharp spot patterns related to bcc-Fe coexisting with bcc-Cr and fee AI crystals, as shown inset of Fig.7(a). The DR has been used in order to determine the crystalline size of the alloy powders in direct way. Figure 7(b) shows the DR and the corresponding SADP of a Fe_.Cr._Nig alloy powder after 86 ks of the MA time. The powder particles have cell-like morphology containing very fine grains with nano-dimensions (about 50 nm or less in diameter). Obviously, the Debye-Scherrer rings of the SADP shows the formation of bcc-FeCrNi solid solution coexisting with unprocessed Cr and/or Ni (fine spots) as shown inset of Fig.7(b). Increasing the MA time leads longer impact shear forces generated by the rods and this causes mechanical defects in the powder particles. Numerous faults and dislocations appear clearly in the particles taken after 22 ks of milling, as shown in Rg.6(b). Moreover, several defects with grain boundary movement are clearly seen near the center of the micrographs. Figure 7(c) snfyvs a high magnification BR of Fe.Cr.„Ni alloy powder after 360 ks of the MA time. The structure of the alloy powder at this stage of milling is fine and a clear halo-pattern is shown in the corresponding SADP, suggesting the existence of an amorphous phase (featureless image). This amorphous phase coexisting with bec FcCrNi solid solution, indicated by the fringe-images and the spot diffraction pattern, as shown in Fig.7(c). At the end of the MA process, the structure of the overall matrix is fine with no dominant facet structure, indicating the formation of an amorphous phase, as presented in Rg.7(d). Moreover, the SADP shows a typical halo-pattem of an amorphous phase in good agreement with the XRD patterns of Rg. 1. The TEM technique allow us lo determine the crystalline size of the Fe Cr Ni alloy powder during the MA process. These results arc in fair agreement compared with the results of XRD using the Scherrer equation [36|, as shown in Fig. 8. The grain size of Fe_.Cr.uNi alloy powder monotonically decreases with increasing the MA time to be almost invisible after 360 ks, i.e. amorphous starting stage. 3.3. Thermal analysis -104- • o 20 nm Fig. 7 TEM micrographs and the corresponding electron diffraction patterns of rod-milled Fe Cr NL alloy powder after (a) 22 ks, i'-,) 86 ks, (c) 360 ks and (d) 1080 ks of the MA time. 1000 i i i • i• 11 i—i i i ini| i i i 11 II fA CalCalcnlatcl d from XRD "\ E 800 .[ I • EstimateEsti d from TEM O 600 E- "53 .5 400 fr OS hi o AmorphizationJ 200 • starts v . $ ^ *, ** • • • » * •"' • • • • * ••n ^ *^ * * * id 10 100 1000 Milling time (ks) Fig. 8 Effect of the MA time on the grain size of rod-milled Fe_ Cr Nig alloy powders. The thermal stability of amorphous Fe74Cr(KNiK alloy powder has been characterized by DSC measurement. Figure 9 displays the typical DSC curves for amorphous Fe_.Cr.„Ni0 alloy powders after 1800 ks of the MA time. During the first heating run (solid line) a single exothermic peak appears at about 1240 K. At the second heating run (dot line), however, this exothermic reaction disappeared. In order to understand the origin of this exothermic reaction, the heated sample after the DSC measurements were analyzed by XRD. The XRD patterns of this sample shows the formation of fcc-Fe_.Cr.uNL alloy powder [see Fig. 10(a)| in good agreement with the #4 lo o XRD patterns of standard austenitic stainless steel, as shown in Fig. 10(b). Moreover, the BFI of this sample shows the formation of large polycrystalline grains with sharp grain boundaries, as shown in Fig. 11. The XRD patterns lead us to attribute the exothermic reaction to amorphous-crystalline phase transformation at crystallization temperature, T . of 1240 K with enthalpy change of crystallization (the area under the exothermic reacliun). AH . of 16.88 kj/niol. 4. DISCUSSION Mechanical alloying (MA) of elemental Fe, Cr and Ni mixture powders under argon gas atmosphere leads to formation of amorphous Fe Cr NL alloy powders. The results have shown that the so-called MA process for producing amorphous FeL.Cr]RNLalloy powders can be classified mainly into three stages. The end-product of each stage varies widely from stage to stage in structure, morphology, and metallography. In this section we shall discuss the mechanism of amorphization for -106- Temperature (C) —*- ,500 600 700 800 900 1000 T r. • -| i i 1 r—" • i 1 1 r lit heating run 0.67 K/s - - • * 2 nd healing run (BUM- line) T, I y I' - - N 3 4 AH, - eo « rod-milled Fe^Cr^M, allov powders a r 1 r • _j_. . . . . i ' 800 900 1000 1100 1200 1300 Temperature (K) —»- Fig. 9 DSC curve of rod-milled Fe Cr NL alloy powders after 1800 ks of the MA time. I 'I I I I I I I I I I I I I I 100 90 80 70 60 50 -#—2 0 (degree) fig. 10 XRD patterns of (a) Fe74CrlhNi8 alloy powders milled for 1800 ks, then heated to 1300 K during DSC measurements and (b) standard sample of austenitic stainless steel. -107- austcnitic stainless steel via MA process through each stage of milling. 4.1. The early stage : During the first few ks of the MA process, e.g. 2 ks, the powder particles of Fe, Cr and Ni blend together without forming any composite particles, as was shown in Figs.l, 3 and 7(a). As a result of cold welding, almost all of the starting material powders consist of assemblages or agglomerations of Fe Cr and Ni to ft».i\ composite particles of a larger diameter particles, as illustrated in Fig. 3(b). This stage of milling can be called also as "negative-milling stage" in which the size of the powder particles increases. The particles at this stage of MA process have well-developed lamellar structure of elemental Fc, Cr and Ni. as was shown in Fig 4(a). 4.2. The intermediate stage: The previous negative-stage of milling is followed by a second stage which is called the intermediate stage or positive stage of milling (22 to 360 ks). Hence, the agglomerated particles are shattered and disintegrated into several particles apparently irregular in shape and size, as was shown in Fig3(c). This disintegration of the powder particles has occurred as result of a continuous shear stress and impact force that has been generated by the rod-powder-rod collision. As the powders are subjected to these external forces, new or fresh active surfaces of the particles appear. During this stage of the MA process, all Bragg-peaks of the bcc-Cr and fcc-Al crystals diffuse completely into the Bragg-peaks of the bec -Fe crystals, indicating the formation of bcc-FeCrNi solid solution, as was displayed in Fig. 1. It is worth noting, that the layered-structure morphology [see Fig.4(b)) as examined by cross-sectional view of the particles, had already disappeared, suggesting the formation of a single phase (bcc-FeCrNi solid solution), as was shown in Fig.4(c). As the milling time increases, 360 ks, the bec-solid solution expands to be 0.2878 nm. In fact, this value is larger than that of pure bcc-Fe crystal (0.2866 nm) suggesting an interstitial solubility of Cr and Ni in Fe, as was illustrated in Fig.2. Further milling leads to mechanical defects of the powders which give rise to several defects in FeCrNi solid solution. These defects are able to change the free energy of the solid solution to less stable phase of amorphous FeCrNi. 4.3. The final stage: In the present study, we shall define the final stage of MA process (360 to 1800 ks) as the homogenization stage in which the solid state reaction takes place homogeneously and a uniform amorphous phase of Fe74Cr.„Ni alloy is obtained. Towards the end of this stage the amorphous phase has XRO patterns with broad and -108- fig. 11 (a) BFI and the (b) corresponding SADP of Fe74Cr|8Ni8 alloy powders milled for 1800 ks, then heated (o 1300 K during DSC measurements. smooth peaks, as was shown in Fig.l. In addition, the amorphous alloy powders crystallize through a single exothermic peak (Fig.9) suggesting that the formed amorphous phase is single phase and homogeneous in composition. This amorphous phase is stable at high temperature and transforms to fcc-austenitic stainless steel at 1300 K, as illustrated in Fig. 10. 5. CONCLUSION A new process for producing austcnitic stainless steel powder at room temperature is reported. The amorphous FeCrNi alloy has been fabricated by milling elemental Fe, Cr and Ni powders in a rv mill under purified argon gas at room temperature. The results have shown that the mode of amorphization of Fe7.Cr,„NLalloy powders via rod milling technique can be classitied into three stages of milling as follows: (1) At the first stage, the elemental powders of Fe, Cr and Ni particles form layered-composite particles of a larger diameter as a result of cold welding. (2) At the second stage, the elemental Cr and Ni powders have been completely diffused into the Fe matrix to form a bcc-FeCrNi solid solution. This solid solution expands with increasing milling time and leads to a saturation value in the lattice parameter, a , of 0.2878 nm after 360 ks of milling. (3) At the final stage of milling, further mechanical defects that resulted from increasing the MA time lead to a transformation bcc-FeCrNi solid solution into the less stable phase of amorphous Fe- Cr NL alloy powders. -109- The crystallization characteristics presented by crystallization temperature, T ., and the enthalpy change of crystallization, AH , arc 1240 K and -16.88 kJ mol" , respectively. One goal of the present study is to offer a powerful and easy technique for formation of stainless steel alloy at room temperature. REFERENCES (1) R.B. Schwarz and W.L. Johnson, Phys. Rev. Lett.. 51. (1983) 415. (2) Y.T. Cheng, W.L. Johnson and M.A. Nicolet, Appl. Phys. Lett., 47 (1985) 800. (3) H. Schroder, K. Samwcrand U. Kostcr, Phys. Rev. Lett., 54(1985) 197. (4) A.Y, Yermakov, Y.Y. Yurchikov and V.A. Barinov, Phys. Met. Metall., 52 (1981) 50. (5) R.L. White, Ph.D Thesis, Stanford University, 1977. (6) C.C. Koch, O.B. Cavin, C.G. McKamey and J.O. Scarbrough, Appl. Phys. Lett., 43 (1983)1017. (7) S.R. Elliott, Physics of Amorphous Materials (Longman Scientific & Technical, (1990) p. 25. (8) J.S. Benjamin, Met. Trans., JA (1970) 2943. (9) I.G. Wright and A. Wilox. Met. Trans.. 5A (1974) 957. (10) J.S. Benjamin and T.E. Volin, Met. Trans., 5A (1974) 1929. (11) J.S. Benjamin, Sci. Am., 40 (1976) 234. (12) G.H. Gessinger, Met. Trans., JA (1976) 1203. (13) J.S. Benjamin and MJ. Bomford, Met Trans., 8A (1977)1301. (14) P.S. Gilman and W.D. Nix, Met. Trans., J2A (1981) 813. (15) J. S. Benjamin and R.D. Schelleng, Met.Trans.. 12A (1981) 1827. (16) K. Suzuki, J. Non-Cryst. Solids. 117/118 (1990) 1. (17) M. Sherif El-Eskandarany, F. Hoh, K. Aoki and K. Suzuki, J. Non-Cryst. Solids, 117/118(1990)729. (18) M. Shcrir R-Rskandarany, K. Aoki and K. Suzuki, Scripla Mclall., 25 (1991) 1695. (19) M. Sherif El-Eskandarany, K. Aoki and K. Suzuki, J. Appl. Phys., 71. (1992) 2924. (20) M. Shcrif El-Eskandarany, K. Sumiyama, K. Aoki and K. Suzuki, J. Mater. Res., 7 (1992)888. (21) M. Sherif El-Eskandarany, K. Sumiyama, K. Aoki and K. Suzuki, Mater. Sci. Forum. 88-90 (1992)801. (22) M. Sherif El-Eskandarany, K. Aoki and K. Suzuki, AppL Phys. Utt., 60 (1992) 1562. (23) M. Shcrif El-Eskandarany, K. Sumiyama, K. Aoki, T. Masumoto and K. Suzuki, J. Mater. Res. (1994) in the press. -110- (24) M. Shcrif El-Eskandarany, K. Aoki and K. Suzuki, J. Less- Common Met., 167 (1990)113. (25) M. Shcrif El-Eskandarany, K. Aoki and K. Suzuki, J. Jp.i. Soc. of Powder and Powder Metal!. (JSPM), 38 (1991) 934. (26) M. Shcrif El-Eskandarany, K. Aoki and K. Suzuki, Forum, 88-90(1992)81. (27) M. Sherif El-Eskandarany. K. Aoki and K. Suzuki, J. Non-Cryst. Solids, 150 (1992)472. (28) M. Shcrif El-Eskandarany, K. Aoki and K. Suzuki, J. Alloys Comp.. 186(1992) 15. (29) M. Sherif El-Eskandarany, K. Aoki and K. Suzuki, Met. Trans., 23 A (1992) 2131. (30) M. Sherir El-Eskandarany, K. Aoki and K. Suzuki, J. Appl. Phys., 72 (1992) 2665. (31) M. Shcrif El-Eskandarany, K. Aoki, H. Itoh and K. Suzuki, J. Less-Common Met., 169(1991)235. (32) M. Sherif El-Eskandarany, K. Aoki and K. Suzuki, J. Alloys Comp.. 177(1991) 229. (33) M. Sherif El-Eskandarany, J. Alloys Comp., 201(1994) 117 (34) M. Sherif El-Eskandarany, in proceeding of the Al-Azhar Engineering Third Internationa! Conference, December 18-21, 1993, Vol. 9, pp. 129. (35) M. Sherif El-Eskandara;i;, Ph.D Thesis, Tohoku University, Sendai-Japan, 1992. (36) A. Guinier, X-ray Diffraction (Freeman, San Francisco, CA, (1963) p.124. -Ill- First International Spring School & Symposium on Advances in Materials Science (SAMS 94) 15-20 March 1994 FORMATION AND CHARACTERIZATION OF NICKEL TITANIUM HYDRIDE ALLOY POWDER M. SHERIF EL-ESKANDARANY' AND H.A. AHMED2 I. Mining and Petroleum Engineering Department, Faculty of Engineering, Al-Azhar University, Nasr City, Cairo, Egypt. 2. Mining, Petroleum and Metallurgical Engineering Department, Faculty of Engineering, Cairo University, Giza, Egypt. ABSTRACT Amorphous nickel titanium hydride alloy powder has been synthesized by milling an cquiatomic mixture of elemental nickel (Ni) and titanium (Ti) powders in a rod mill under hydrogen (H) gas atmosphere at room temperature. The mechanically alloyed powders have been characterized by means of x-ray diffraction, differential scanning caiorimetry, optical microscope, scanning electron, and transmission electron microscopes and chemical analysis. After 11 ks of the rod milling time, the coarse powder particles of Ni and Ti were disintegrated into several particles that have fresh surfaces. These fresh or new surfaces are very active and able to absorb hydrogen gas, so that hcp-Ti reacts completely with the hydrogen gas to form fcc-TiH, with grain sizes of about 60 nm in diameter. No phase transformation has been observed in fcc-Ni. After 43 ks of rod milling time, the formed fcc-TiH, diffuses into Ni matrix to form fec- NiTiH3 solid solution with average grain size of 10 nm in diameter. Further milling (360 to 720 ks) creates mechanical defects that lead to raise the free energy of fcc-NiTiH3 solid solution to less stable phase of amorphous NiTiH3 alloy powders. The end product of the amorphous NiTiH3 alloy powder has spherical like morphology with average size of about 3 fim in diameter. After 720 ks of the rod milling time, the crystallization temperature and the enthalpy change of crystallization for amorphous NiTiH3 alloy are 774 K and -7.88 kJ/mol, respectively. Key Words: Mechanical alloying, Rod milling, Solid state amorphization reaction, Powder metallurgy, Structure, Morphology, Caiorimetry, Metal Hydride -113- 1. INTRODUCTION Mechanical alloying (MA) process via ball milling technique has bten successfully used for formation of composite powder particles with intimately dispersed, uniform internal structure (1). In 1983, Koch etal. (2) gave the first example of formation of Ni^jNbjQ amorphous alloy powder using a high-energetic ball mill. It is believed that the mechanism of amorphization by mechanical alloying is similar to that for amorphization by solid state amorphization reaction, SSAR, (3). Comparable to MA process, amorphous alloys can be prepared by grinding the crystalline compound, that is to say, mechanical grinding, MG (4) or mechanical disordering, MD (5). Since then, there have been several examples of metallic amorphous alloys fabricated by MA process using ball milling (6-14) and/or rod milling (15-23) techniques. It is worth to say that the MA process is used mainly as a successful method for producing amorphous alloys that arc difficult or impossible to be obtained by conventional melting techniques, e.g. Al-Ta (10) and Al-Nb(24,25) binary systems. It has been also used for the amorphization of the systems exhibiting positive heat of mixing (26,27). Another attractive novel technique for formation of metal nitrides alloy powders by milling the elemental powders in a nitrogen gas flow at room temperature, has been first reported by H-Eskandarany el al. (28). This unique technique that called reactive ball milling (RBM) gave another application of the ball milling method for preparing several metal nitrides alloys (29-32). In the present study, an amorphous of NiTiH3 alloy has been synthesized by milling an cquiatomic mixture of elemental Ni and Ti powders in a rod mill under a hydrogen gas atmosphere. The progress of crystal-to-amorphous phase transformation in NiTiH3 alloy powders has been followed by x-ray diffraction and transmission electron microscopy. Moreover, scanning electron microscopy has been used to understand the morphological metallographica! characteristics of the alloy powders. In addition, the crystallization characteristics of the formed amorphous phase are examined by differential scanning calorimeiry The mechanism of hydrogenation reaction for the formation of a hydrogenous amorphous NiTi alloy via rod milling technique is discussed. One goal of the present study is to offer a unique process for formation of metal hydrides at room temperature by a simple technique. Finally, this is a systematic study of a system that has no reported thermodynamic data. -114- 2. EXPERIMENTAL High purity elemental Ni (99.9 %, 70 urn) and Ti (99.5 %, 50 p.m) powders, and purified hydrogen and argon gas (H,0 and Q, < 10 ppm) have been used. The powders were mixed in a glove bag under a purified argon atmosphere to give the desired average composition of Ni^Ti^ The mixed powders were then charged and sealed in a cylindrical stainless steel shell (SUS 304, 120 mm in diameter) together with stainless steel rods (SUS 304, 10 mm in diameter). The rod-to-powder weight ratio was controlled to be about 36:1. The inlet of the shell was connected with a rotary pump and evacuated for about 4 Its. Then a flow of hydrogen gas was passed carefully into the rod mill through a plastic pipe. Once the rod mill was filled with 1 atm of hydrogen, the inlet of the vial was closed and the reactive rod-milling process was carried out at ambient temperature by mounting the rod mill on a rotator at the rate of 1.4 s" . The reactive rod milling was interrupted at selected intervals and * small amount of the rod- milled powder was taken out from the vial in the glove bag. The alloy powders were characterized by x-ray diffraction (XRD) with CuKct, scanning electron microscopy (SEM) operated at 25 kV and transmission electron microscopy (TEM) using a 200 kV microscope and differential scanning calorimetry (DSC) under an argon gas flow at heating rate of 0.67 Ks . The induction coupled plasma (ICP) emission method was used to analyze the contents of Ni and Ti, and the degree of Fe contamination in the milled powder. The oxygen and the hydrogen contents in the alloy powders were detected by the helium carrier fusion-thermal conductivity method. After 720 ks of milling, the iron and oxygen contents in the alloys were determined to be less than 0.60 and 0.90 at%. respectively. 3. RESULTS 3.1. Morphology and metallography The morphology and metallography changes with the rod milling time of Ni50Ti5Q alloy powders have been followed by SEM technique. Figure 1 displays SEM micrographs of Ni^Ti^, alloy powders after selected MA times. After 22 ks of the MA time, the cross-sectional view of a polished and etched agglomerated particle shows a typical lamellar structure of Ni (light gray matrix) and TiH2 (dark gray veins). Increasing the MA time (173 ks) leads to decrease the average size of the powder particles from nearly 600 |im (Fig. l(a)l to almost 40 (im in diameter, as presented in Fig. 1(b). After this stage of milling, the layered structure morphology is hardly observed (Fig. I(b)| suggesting th* formation of single phase. At the end of the MA time, the particles have been dramatically decreased in size to get an average homogeneous size of about 3 um and the shape of these particW«, changed from irregular flake-like morphology to uniform spherical-like morphology, as shown in Fig. 1(c). -115- XRD has been used in order to understand the total structure of the milled powder after several stages of the MA time. Figure 2 illustrates the XRO patterns of rod milled NisoTi'so alloy powders as a function of the MA lime. At the initial stage (0 ks) of milling, the starting material consists of polycrystalline of fcc-Ni and hcp-Ti. After 11 ks of the MA time, the Bragg-peaks corresponding to hcp-Ti were surprisingly disappeared and a new phase of fcc-TiH2 alloy is detected, indicating a complete reaction between the milling atmosphere (hydrogen) and hcp-Ti. The lattice parameter, ag, of TiH, phase was calculated to be 0.2871 nm. After 43 ks of the MA time, the Bragg-peaks of fcc-TiH, alloy have become broader and their.intensities are sharply decreased. Moreover, the positions of the Bragg-peaks for fcc-Ni crystal [e.g. Ni (200)] have shifted remarkably to the low angle's side, indicating the formation of partial fcc- NiTiHj solid solution, as shown in Fig.2. The ag of the formed solid solution expands with increasing the MA time to get a maximum of 03549 nm after 173 ks. It is worth noting that at this stage of milling, all the Bragg-peaks for fee-'UK-, crystal have completely disappeared. This suggests the -116- formation of single phase of supersaturated NiTiHj solid soluti<>*. alloy powder. Towards the end of the MA time (360 Its), these peaks become broader with diffuse haloes, indicating that NiTiHj solid solution has been completely transformed to a homogeneous amorphous phase. We should emphasis that this amorphous phase has not been changed to any other phase(s) even for the samples nv'lcd for more longer time as long as 720 and/or 1440 ks, as illustrated in Fig. 2. 3.2.2. TEM TEM analysis has been used to understand the local structure of the milled powders at several stages of the MA process. Figure 3 displays the bright field image, • I • • I • I I I I T I I I I I J I I • 11 • • ' • I ' ' ' ' I • • I'l"" I" " Oks §= CO zP •*— 2 e II S - t-i 11 ks a Tilli (111) u •a* 43 ks < CO E 0) l •. •_> - • 80 70 60 50 40 30 -*— 2 0 (degree) Fig. 2 Typical XRD patterns of equiatomic NiTi powders milled in hydrogen atmosphere after selected MA times. -117- Rg. 5 High magnification TEM micrograph for N1T1H3 all»»;' powders after 22 ks of the MA time. BFI (a) and the corresponding selected area diffraction patterns, SADP (b) of rod milled NiTiHj alloy powders after 4 ks of the MA time. Obviously, the powder particles are heavily faulted containing polycrystalline of Ni and Ti with an average grain size of about 300 nm in diameter. Figure 4 shows the BFI and SADP of rod milled NiTiHj alloy powders after 11 ks of the MA time. Remarkably, hcp-Ti powders react with hydrogen and completely transformed to fcc-TiH2 (zone 1) as shown in Fig. 5(b). Contrary to this, no phase' transformation in fcc-Ni crystals could be detected, as illustrated in Fig. 5(c). The powder particles are almost equiaxed with average grain size of about 60 nm in diameter. Fgure 5 shows a high magnification BR (a) and the corresponding SADPs (b) and (c) of rod milled NiTiH3 alloy powders after 22 ks of the MA time. The shown -118- Fig. 3 TEM observation for NiTiH3 alloy powders after 4 ks of the MA time. Fig. 4 The BFI and the corresponding SADP of NiTiHj alloy powders after 11 Its of the MA time. 86 ks. : l l^r^?' » •2. 4 J. . 4* 41, ). «>" -(- .-• ^ A kJil.'.l Rg. 6 The BR and DFI are shown logether with SADP of NiTiH3 alloy powders after 86 ks of the MA time. -120- NT" £.« N> 173 ksl EB1M Hg. 7 The DFI is shown together with SADP or NiTiHj alio) powders after 173 ks of the MA time. X z 1 a. •oc a. a < 2 o r- ca H ? a -122- fringe image corresponds lo fcc-TiHi, suggesting a solid stair, reaction between hcp-Ti powders and hydrogen, as illustrated in Fig.5(c). Contrary to this, no phase transformation in fcc-Ni crystals could be delected, as illustrated by the matrix shown in Tig. 5(a) and the SADP of Fig. 5(b). Figure 6 shows (a) BFI and (b) dark Held image, DFl, of Ni'l'iH, alloy powders after 86 ks of the MA time. The corresponding SADPs arc displayed inset of Fig. 6. Two regions are indicated in the BFI (I and II) and one region is indicated in the DFl (I/O. The overall matrix hi*, cell-like morphology of crystalline fcc-Ni |Fig. 6(c)| and fee-Til-^ IFig. 6(e)|. The morphology of the region II, however, has i fine structure and thr corresponding SAUP indicates ihe rornidfiou of in amorphous phase coexisting wjlh sharp spot diffraction pattern, as shown in Fig. Fig. 6(d). The origin of these spots pattern originating from unprocessed fcc-Ni crystal, as shown in Fig. 6(b). Figure 7 displays the DFl and Ihe corresponding SADP of NiTiH3 alloy powders after 173 ks of Ihe MA lime. The alloy panicles have fine lens or grain-like morphology wilh average size of 10 nm in diameter. The SADP shows Ihe formation of fee N1T1H3 solid solution, as shown in Fig. 7(b). The BFI and Ihe corresponding SADP of a near edge particle taken after 720 ks of Ihe MA time, are shown in Fig. 8. Overall, the alloy powder particles appear to have homogeneous fine structure with no dominant facet structure. Moreover, (he SADP shows a typical halo-pattern of an amorphous phase in an excellent agreement with the XRD patterns of Fig. 2. The crystalline size of NiTiHj alloy powders using TEM technique is plotted against the MA time in Fig. 9. The size of these grains monotonically decreases v.'ith increasing the MA time to be almost invisible after 259 ks, i.e. arnorphiution starting stage. .?_?. Thermal analysis Parallel to the investigations by SEM, XRD and TEM, the progress of the solid stale amorphization reaction has been characterized by DSC measurements. Figure 10 displays the DSC curves for NiTiH, alloy powders at the several stages of milling. For the same sample, the solid line presents the first heating run, whereas the dashed line show the second heating run. The exothermic reactions of the first heating run can be identified as amorphous-to-cryslalline phase transformation (crystallization reaction) of the amorphous phase in the alloy powders. The peak temperatures for the first heating run that present the crystallization temperature (TK) were plotted Fig. 11 as a function of the MA time. Meanwhile, the data from the second scan were subtracted from the data of the first scan and plotted in Fig. 12 in terms of enthalpy change of crystallization (AH,). After 43 les of the MA time, the first healing run reveals a single broad -123- 90 T-r'r-i ri>| '•' I"11 "»' TIITI. • r' I I i Jl M; X T H .3 - T -I Amorphous • I r i mil i i i i i ml i tji0 nff II 111 lioi Milling time (ks) Fig. 9 Grain size distribuiion of NiTiHj alloy powders as a function of the MA times. 1 ' 0.67 K/s I 1J i i ' . • i . i 400 500 600 700 800 900 Temperature (K) Fig. 10 DSC curves for NiTiHj alloy powders after rented MA times. -124- 800 i iii| i 700 J * / /• X 600 y c l J 500 '— - - • *"' J 10 100 1000 Milling time (ks) Fig. 11 Crystallization temperature, Tx, for NiTiH3 alloy powders as a function of the MA time. u 1- f-l-I'TTTTf- 1 . "o 2., ; \ X X < - 10 100 1000 Milling time (ks) Fig. 12 Enthalpy change of crystallization, AT,, for NiTiH3 alloy powders as a function of the MA time. -125- exothermic peak appears at about 560 K. This peak temperature is shifted somewhat to the elevated temperature (650 K) and becomes pronounced after 173 ks of the MA time. This suggests a volume fraction increasing of the amorphous phase in the alloyed powders. Towards the end of the MA time (360-1440 ks) the exothermic peak becomes more pronounced and very sharp. as displayed in Fig. 10. Figure 11 illustrates the Tx of mechanically alloyed NiTiHj powders as a function of the MA time. At the first stage of milling (43-360 ks). the value of T, increases sharply, indicating a continuously dramatic change of the composition of the amorphous phase and approaches a saturation value of 778 K towards the end of the MA time (720- 1440 ks). The AHX (the area under the crystallization peak) being a measure for the quantity of the amorphous phase in mechanically alloyed NiTiH3 powders, is shown in Fig. 12. During the first stage of milling the value of AHX decreases dramatically, suggesting a large increase in the volume fraction of the amorphous phase. At the end of the MA time, AHX remains nearly constant at a value of -7.6 kj/niol. This confirms the formation of homogeneous amorphous alloy. 4. DISCUSSION Static hydrogen-induced amorphization has been tradi'icnally used for formation of several metallic amorphous hydride alloys (33-34). The present study shows that MA of an equiatomic mixture of Ni and Ti powders via reactive rod milling under hydrogen gas atmosphere at room temperature leads to the formation of homogeneous amorphous NiTiHj alloy powders. Moreover, the rod milling for preparing amorphous NiTiH3 alloy powders can be classified into three stage of milling, i.e. early, intermediate and final stages. The structure of the end product of the alloy powders differs widely from stage to stage. In this section the role of amorphization of NiTi alloy via dynamic- assisted solid state hydrogenation reaction using the rod milling technique is discussed in structure, morphology and calorimetry points of view. Figure 13 will be used to summarize the mechanism of solid state hydration reaction by way of rod milling technique. 4.1. The early stage of milling: This stage of rod milling (0-43 ks) refers to the behavior of the powder particles of Ni and Ti through the hydration reaction at the first few kilnseconds of the MA time under hydrogen atmosphere. At this stage, the starting elemental powder particles of Ni and Ti [Fig. 13(a)) are mainly subjected to dramatic continuous shear stress that generated by the rod as milling media (Fig. 13(b)|. The particles are disintegrated into several particles and very clean or fresh active surfaces of the powder particles are created, as illustrated in Fig. 13(c). Hence, the reactive hydrogen gas that initially -126- Tlptrikt* FmJi inrfac* (b) of Ni fmrtkk Starting stage, mixture of Hod-powdei-rod Early stage of milling, formation elemental Ni and Ti powder. action. of fresh surfaces pfNi and Ti powders. Solid Comport f e particle ol \ NlmdTUIS (0 Solid state hydration reaction, Formation of layered- Intermediate stage of formation ofTitia. structure powder particles. milling, formation of fcc-Nttilh solid solution containing amorphous phase alloy. (X) Final stage of milling formation of uniform and homogtntous amorphous ofNifflh alloy powder. Fig. 13 Proposed flowsheet diagram illustrates the mechanism of solid state hydration and amorphization of NffiHj alloy powders by mechanical alloying. presented in the rod mill was gettered and absorbed completely by the first atomically clean surfaces of Ti powders to form fcc-TiH2 alloy, as shown in Fig. 13(d). Compared with Ni, Ti as 4f element has high ability of for taking up hydrogen to form metal hydride. Since the hcp-Ti has been transformed completely to fec-TiHi phase, the MA process was changed from a solid state hydration reaction between the elemental Ni and Ti powders to a solid state reaction between fcc-Ni and fcc-TiH, powders. Increasing the MA time leads to the formation of composite lamellar-structure |see Fig. I(a)| powder particles containing Ni and TiH2 phases, as displayed in Fig. 13(e). -127- 4.2. The intermediate stage of milling: At the intermediate stage of milling (43-360 ks) the Jarmed phase of fcc-TiH; diffuses completely into the unreacted phase of fcc-Ni to form supersaturated fcc- NiTiHj solid solution. As the MA time increases the a<, of NiTiH3 solid solution expands to be 03549 nm after 173 of the MA time. It is worth noting that this value is larger than that of pure fcc-Ni crystal (03524 nm) indicating an interstitial solubility of TiH2 in Ni. It is worth noting that this solid solution containing an amorphous phase of NiTiH3 alloy, suggested by the crystallization reaction peaks of the samples during the DSC measurements (see Fig. 10). In fact, the amorphous phase of this is rather heterogeneous, indicated by the broadness of the exothermic reaction and the large differences of Tx and AH, values. The local structure examined by TEM shows that the powders differ widely in the structure from particle to particle and within the particle, as was shown in Fig 6. Towards the end of this stage, the alloy powder particles have fine nano dimension grains, as presented in Figs. 7 and 9. 43. The final stage of milling: The final stage of milling (360-720 ks) refers to the amorphization process that took place in the fcc-NiTiHj solid solution to form an amorphous NiTiHj alloy powders (Fig. 13(g)). This amorphization process occurs as a result of several defects that able to change the free energy of the solid solution phase to a less stable phase (amorphous NiTiH3). The formed amorphous phase is very homogeneous, suggested by a halo diffused pattern without existence of any other phases (Figs. 2 and 8) and consists of very fine powder particles {about 3 um in diameter) in a narrow particle size distribution, as displayed in Fig. 1(c). Moreover, the internal structure is fine (Fig. 8) without nanocrystalline phases (Fig. 9). The thermal stability of this amorphous phase indexed by the values of Tx, and AHX has been increased, as presented in Figs. 11 and 12 respectively. 5. CONCLUSION Low energetic rod mill has been used via mechanical alloying process for preparing an amorphous nickel titanium hydride alloy powders. The milling process has been achieved in the presence of purified hydrogen gas at room temperature. It has been the goal of the present study to provide a systematic evolution of the mode of amorphization and crystallization for NiTiH-, alloy powders. The present work allows the following interpretations. 1. After few kiloseconds (11 ks) of the MA time, the coarse powder particles of Ni and Ti were disintegrated into several particles which have new or fresh surfaces that are able to absorb hydrogen gas, so that hcp-Ti reacts completely with hydrogen gas (milling atmosphere) to form fcc-TiH2 with grain sizes of about 60 nm in diameter and a -128- lattice parameter, a^ of 0.2871 nm. Contrary (o this, no phase transformation could be observed in fcc-Ni. 2. After 43 ks of the MA lime, the formed fcc-TiH, diffuses into Ni matrix to form fcc-NiTiH-, solid solution with grain size less than 10 nm in diameter. This solid solution expands with increasing the rod milling time (173 ks) to ff-x a saturation value mag of 03549 nm. 3. Further milling (360 to 720 ks) creates mechanical defects such as point defects and dislocations that lead to raise the free energy of fcc-NiTiH3 solid solution to less stable phase of amorphous N1T1H3 alloy powders. 4. The crystallization temperature, TK, and Ihe heal of Ihe enthalpy change of crystallization, AHt, of amorphous NiTiH3 alloy powders milled for 720 ks were determined 10 be 778 K and -7.6 kJ/mol, respectively. REFERENCES I.BENJAMIN.J.S. Met.Trans.A. 1.2943(1970). 2. KOCH, C. C. CAVIN, O. B., MCKAMEY. C. G. and SCARBOIJROh. J .0.. Appl. Phyj. Lett. 43, 1017(1983). 3. SCHWARZ. R.B. and JOHNSON. W. 1... Ptiys. Rev. 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Less- Common Met, 167,113(1990). 16. EL-ESKANDARANY. M. Sherif. AOKI, K. and SUZUKI, K., 1 .Appl. Phys„ 72. 2665(1992). 17. EL-ESKANDARANY, M. Sherif. AOK1, K. and SUZUKI. K„ Materials Science Forom. 88-90,81 (1992). 18. EL-ESKANDARANY, M. Sherif. AOKI. K. and SUZUKI. K. Met. Trans A 23. 2131(1992). 19. EL-ESKANDARANY. M. Sherir. AOKI. K. and SUZUKI. V.. J. Jpn. Soc. of Powder and Powder Metall. (JSPM). 3D, 59 (1992). 20. EL-ESKANDARANY, M. Sherif. AOKI. K. and SUZUKI, K„ J. Non-Cryst. Solids, 150.472^1992) 21. EL-ESKANDARANY, M. Shcrif, AOKI, K. and SUZUKI. K. J. Appl. Phys.. 71, 2924(1992). -129- 22. EL-ESKANDARANY. M. Sherif. SUMIYAMA. K.. AOKI. K. and SUZUKI. K., J. Jpn. Soc. or Powder and Powder Mclall. (JSPM). 39. a»6(l992). 23. EL-ESKANDARANY. M. Sherif. AOKI, K.. MASUMOTO. T. and SUZUKI. K., Journal of Alloys and Compounds (1993) in the press. 24.HELLSTERN.EandSCHULTZ,L.J Appl. Phys.. 63.1408(1988). 25. EL-ESKANDARANY. M. Sherif. AOKI, K. and SUZUKI. K.. 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Itiys., 58. 4573(1985). -130- First International Spring School & Symposium in Advances in Material Science (SAMS 94) Cairo, 15-20 March 1994 REACTIVE POWDER METALLURGY FOR NIAI-BASE ALLOY AND COMPOSITES K.E. Mohamed*, D. Stover, H.P. Buchkremer • Institute for Applied Materials Research (IAW) Research Center Julich (KFA), Germany *Guest scientist from the Atomic Energy Authority, Cairo, Egypt. Abstract Intermetallic compounds could be the basis for future high temperature structural components. The present work is focused on the use of elemental powders for the fabrication of NiAl-composite, NiAl-Cr as well as NiAl-Cr matrix composite. The primary interest is to study the sintering process during fabrication of dense materials from the elemental form using reactive sintering and reactive hot isostatic pressing processes. Using reactive sintering at 700°C it was possible to get dense NiAl with 97% of theoretical density. The addition of Y203 lowers the amount of densification and with 15V% Y203 a relative density of 80% was obtained after sintering for 15 minutes. Attempts to fabricate the NiAl-Cr eutectic alloy with and without Y203 by reactive sintering are described. At low reactive sintering temperature (700°C), the pressure of Cr and Y2Oo decreased the volume fraction of the liquid formed in the compact, and inhibited its homogeneous spreading leading to swelling and consequently low sinter densities. A shift was made to reactive hot isostatic pressing for higher densification. The paper describes the effect of some processing parameters and the evaluation of the sinter/hipped materials by density measurements, microhardness, EDX analysis, metallography and dilatometric measurements. -131- Introduction In recent years intermetallic phases have become a subject of intensive research with the aim of developing new materials on the basis of these phases, which can be used in a wide range of high temperature applications. Particular attention has been paid to the aluminides especially these based on Ni, Ti and Ta. The intermetallic compound NiAl is one of the most promising candidates due to its low density, high melting point and excellent corrosion and oxidation properties. At present, however, the use of NiAl is limited by its low room temperature toughness and high creep rates. There are various possibilities for improving the high temperature strength' ' which inlcude alloying with eutectic forming elements' ' as well as second phase dispersion 13,4] strengthening . Fine TiB2 particles of 30% (l-3^m) were found to increase the compressive flow strength of stoichiometric NiAl manufactured by the XD process from 25 MPa to 95 MPa at 1027°Ct5, 61 . Also, the addition of 2.7% TiB2 to NiAl fabricated by rapid solidification increased the tensile yield stress from 85 MPa to 208 MPa at 760°Ct7]. An increase in the compressive yield strength from 275 MPa to 385 at 700°C due to incorporation of 20V% TiB2 into NiAl fabricated by reactive hot isostatic pressing (RHIP) was reported'4' . Directional solidification was also used to produce NiAl reinforced with Cr and Mo fibers'81. Powder metallurgy processes provide a fabrication route of the intermetallic compounds which has definite advantages over the casting and forging technology. Using the powder techniques the segregation processes are minimized especially for small parts. Aslo, powder processing provides a promising route for fabrication of intermetallic matrix composites' ' '. Many powder metallurgy techniques have been used to fabricate aluminide intermetallies. These include reactive sintering, reactive hot isostatic pressing (RHIP), hot isostatic pressing (HIP), vacuum hot pressing, powder extrusion, -132- explosive compaction, plasma spraying and powder injection molding. In particular, reactive processing has shown great potential for forming nickel aluminides. It has the advantage of using mixed elemental powders, thus avoiding the expenses and technological difficulties encountered in processing prealloyed powders. Reactive sintering involves reacting elemental powders to form the desired compound^ • . Upon heating a powder mixture to the lowest liquidus temperature of the system, a transient liqxiid forms which spreads through the compoenent, consuming the elemental powders and precipitating a solid intermetallic compound'1 . The driving force for this process is the thermodynamic stability of the aluminide intermetallic. For the Ni-Al-system, the first liquid is an eutectic liquid which forms at the interface between contacting particles. The temperature of initiation is about 650 .C[14] and depends on the particle size, powder composition and heating rate' '. The reaction enthalpy for the formation of NiAl is^141; Ni + Al -» NiAl H29g = -118.4K cal/mole. Because of the exothermic nature of the reaction, heat is released which markedly increases the temperature of the sintered compact. Reactive sintering is spontaneous once the liquid forms and the reaction starts. The transient liquid formed in the system provides a capillary force on the structure which leads to rapid densificatiom1 '. The great potential of reactive sintering as a fabrication route for intermetallic matrix composites has been emphasized in several recent reports^ ' " ' • In these investigations the effect of various processing parameters (sintering temperature, powder particle size, green density, composition, heating rate and sintering atmosphere) has been studied and they offer a basis for preselection of suitable reactive sintering conditions. Because of rapid spreading of liquid during this process, pore formation is common and dimensional control often proves difficult if an excess of liquid is formed. The dimensional control difficulties were overcome by addition of prealloyed NiAl powder and/or TiB2 particles^ ' ' . The exact amount of prealloyed powder depends on the compact mass and the powder particle size, but generally ranges from 15 to 25%. Alloying of NiAl by third element addition during reactive sintering was tried^ ' but only in the low level concentration range; 1% Ti, 1% Nb or 5% Cr. the present investigation is an attempt to utilize the heat liberated during the exothermic reaction between Ni and Al in developing high Cr NiAl-alloy from elemental powders. Fine Y^Oo powder was added to NiAl or NiAlCr powder mixtures as a strengthening and dimensional control agent. Reactive sintering, reactive hot isostatic pressing and high temperature liquid phase sintering processes are tried in the present work. Measures for the success in this respect are sintered homogeneous microstructures, low final porosity suitable for full densification by containerless HIPing and good shape control. Experimental Approach The pure elemental powders used in the present study were INCO Ni3-7|im; Al powder of 20 nm size and <25 Jim Cr powder. Prealloyed NiAl atomized and milled < 45 um was used for dilution during the reactive synthesis. Y20„ was added in the form of powder agglomerates < 10 |im diameter. Unless specified otherwise, the compositions are given in weight percent. Appropriate elemental powder mixtures corresponding to NiAl and NiAlCr matrices with 15% NiAl and the required amount of reinforcing Y203 phase are mixed for 6 hours in turbular mixer. Cylenderical compacts were made (6 gm.) by cold isostatic pressing (CIP) of the mixed powders in a polyurethane bag at 350-400 MPa for 15 minutes. The resulting green compacts have density from 65-70% of theoretical. Reactive sintering of the CIPed compacts was -134- performed in vaccum 10 Pa). The process was conducted at 700°C for 15 minutes with a heating rate of 15°C/minute. For RHIPing, the green compacts were canned in thin walled 316 stainless steel tubes, degassed and electron beam welded under vacuum. The RHIP experiments were carried out at 1200°C for 3 hours using a pressure of 172 MPa. The pressure was applied at a rate of 3.5 MPa/min. with a heating rate of 15°C/minute. Dilatometric measurements were carried out using a dilatometer facility installed on a HIPing machine. Dilatometry was performed to monitor the reaction temperature and the specimen length changes with time and temperature during reactive sintering. Postdensification investigations consisted of density and microhardness measurements as well as optical metallography. Electron microprobe analysis was conducted for composition and phase identification. Results and Discussion Reactive Sintering In spite of the addition of 15% prealloyed NiAl powder to the powder mixture corresponding to NiAl stoichiometry (69.5 Ni-30.5 Al), reactive sintered compacts which did not contain Y203 have completely lost their cylenderical shape (Fig. 1) due to the excess amount of liquid formed in this case. The degree of distortion decreased with increasing Y203 content and with > 10 v % Y203 good shape control was attained. Although YgOg provides shape control of the compacts, it hindered densification. Reactive sintered density of monolithic NiAl was found to be 97% of theoretical. With Y203 addition of 15 V% the relative reactive sintered density decreased to 82% (Fig. 2). Etched microstructure (Fig. 3) shows that the addition of Y203 caused grain size refining of reactive sintered NiAl. The grain size decreased from 45 jm -135- B 10 mm Figure i: Reactive sintered NiAl Compac en NiAl. ts with 15 % prealloyed a / b/ c: 0, 5, 10 V % Y2o3 0% 5% 10% 15% Amount of Y203(V%) ^VHN | rel. density Fig. 2. Microhardness and density of reactive sintered NiAl with different amounts of Y2 O 3• Fig. 3. Microstructure of reactive sintered NiAl a) 0% Y20v b) 15 V% Y^Og. -137- from monolithic NiAl to 20 urn for NiAl - 15 V% Y203 . Figure 3 also shows the porosity associated with the presence of Y2Og. Electron microprobe analysis showed that the light and dark-etched grains differ in Al content by about 2%. Figure 2 reveals that the microhardness is increased by the addition of Y203 . The measured value of 257 VHN for monolithic NiAl agrees well with values of 275 VHN reported in previous works14, u\ Following the analysis made by Alman and Stolotf ' for TiB2 reinforced NiAl, a large portion of the strengthening observed in reactively sintered NiAl-YgOg composite can be attributed to grain size refining. In addition, contribution from particle matrix interaction should be taken into consideration. For reactive sintering of NiAlCr alloy, elemental powder mixture containing 39% Ni; 23% Al and 38% Cr (corresponding to NIAl-Cr eutectic system) was mixed with different amounts of Y203 ranging from 0-15 V%. In some cases 15% prealloyed NiAl powder was also added. Sintered compacts containing 15% prealloyed NiAl showed a high degree of porosity which increased with increasing Y2Og content and the density was less than the green one (Fig. 4). Compacts without prealloyed NiAl, in the absence of Y203, have lower porosity (Fig. 5-a). In the presence of Y203, however, the removal of prealloyed NiAl did not lead to remarkable changes in density (Fig. 6). For both cases, the densified areas have a microhardness which is higher than that of NiAl manufactured by the same process (Figs. 5, 6). Small additions of 3,55 or 5 V% Y203 caused a marked increase in the microhardness of NiAlCr material. In transient liquid-phase sintering, the liquid quantity, distribution and duration dictate the final density*21' 23l For reactive sintering, the presence of different unreacting materials cause discontinuities in the advancing liquid interface and a decrease in its volume fraction in the -138- 0% 5% 10% 15% AmounlofY203(V%) ESS VHN [ rel. density Fig. 4. Microhardness and density of reactive sintered NiAl 38% Cr containg 15% prealloyed NiAl. -139- Fig.5. Photomicrographs of reactive sintered NiAl-38% Cr without prealloyed NiAl. a) 0% Y20_ b) 10 V% Y 0 . 0.00% 3.35% to.00% Amount of Y203jtf%) jESSVHN BB f.Ldensity . Fig.6. Microhardness and density of reactive sintered NiAl-38% Cr without prealloyed NiAl. -140- compact. As the volume fraction of Y90„ is increased or Cr is introduced, the degree of interconnectivity of liquid phase is decreased. Thus, no long range capillary action is possible and desnification is inhibited. Also, if the solubilities between formed NiAl-phase and the unreacting materials are unbalanced, swelling is expected. As tthe concentration of Cr in the NiAl-38% Cr is beyond the solubility limit for Cr in NiAl(24>, this may explain the additional effect of Cr on pore formation during reactive sintering of NiAl-38% Cr alloy. Another important factor is the thermal expansion mismatch parameter between NiAl and the foreign unreacting Y203 and Cr which can also play a role in pore formation during reactive sintering. Reactive sinterability results of NiAl and NiAl-38% Cr with Y2Og were further substantiated by dilatometric measurements. Compacts of different composition were reactive sintered in dilatometer facility. The results of sample length changes and the reaction temeprature monitored are given in Table 1. For NiAl with 10 V% Y2Og there is a dimensional decrease of 24, 42% which means shrinkage and densification. The measured density is about 90% of theoretical. The addition of 38% Cr causes drastic dimensional increase in the presence of Y203. Compared to NIA1, at the same content of prealloyed NiAl and Y203, Cr caused dimensional increase of 15,87% instead of 24,42% dimensional decrease. This caused swelling of the compact and the appearance of surface cracks. In fact, Y203 is highly responsible for the swelling observed in the NiAl-38% Cr compacts. The addition of 3,33V% Y203 in the absence of prealloyed NiAl caused 10,84% dimensional increase, while the addition of 15% prealloyed NiAl in the absence of Y203 caused 12% increase. It should be emphasized that the addition of 38%Cr in the absence of prealloyed NiAl and Y203 causes only 2,38% dimensional increase. In this case, the -141- Table 1: Results of Dilatometric Measurements during Reactive Sintering at 700* C for 15 Minutes With 15'C/min. Heating Rate. f mmcriils conieni of contenw, reiciion 1 prc.lloycd NiAJ # ! Tciapenturc W ( *M »s 10 •24,42 j 480 I NiAJ.JSCr »» 0 112.00 ! 460 j NiAl-SSQ ;< 30 -15.S7 j 4S0 j NiAl-SSC: o 0 +2.38 1 470 1 NiAJ-SSCr 0 3.33 *10.« j 470 J NiAloSCr 0 30.6 4)8.30 ; 545" • h:i;:r.f me n *joui y C .r.:n. e 2: Electron Microprobe Analysis of some NiAl-Cr Alloys processed By Reactive Synthesis from Elemental Powders NiAltich phase Cf-fich phAse pnwJci riimpoMlimi CClllMllul.llioil pjf*C<» Ni Al Cr Ni Al O 3<» Ni - 2'. Al - 3K O OP 4 icnriivr hip I5V%Y,0:, (I2MVC. mMl'n. 3lnM 69.S2 25.6K '4.5 4.52 R.92 Kft.56 •endive Miuciing 37.6 Ni-24.4 AI-3SO ai 1425'C. 2 hi* 6.1.82 24.38 .11.79 H.32 7.61 W.07 Cii' • rcaclivr hip 55 Ni • 25 Al. 20 O (1200* C. I72MP.1. 3 his) 71.59 22.X4 5.57 49.09 16.33 34.5K reactive sintering M.56Ni-23.4AI-25Cr ml425*C.2hr* • 66.54 23.75 9.71 31.10 6.24 62.6 59 Ni • 25 Al - 3.S Cr Mulcting al picnllnvcd niomucd pmvilcr 1415*C+ hip 70.25 26.0S 3.66 2.65 4.66 92.6'' surface cracks disappear and the density is about 68% of theoretical. These results show that process parameters like type and content of unreacting materials affect the amount of liquid phase formed and its distribution in the microstructure, thus leading to different degrees of densification after reactive sintering. For a heating rate of 15"C /min, the reaction temperature ranges 460 - 480°C independently of Cr addition or Y2Og content (Table 1). The reaction temperature is however sensitive to the heating rate and was increased to 545°C for a heating rate of 5°C/minute. This mean? that the exotherm occurs prior to liquid formation from the first eutectic in the Ni-Al system (640°C) or from melting of Al (660°C). In this case, the liquid forms due to self heating of the compact. Similar observation was reported^ ' for reactive sintering of NinAl where dilatometry and differential thermal analysis indicated that the reaction occurs prior to 600°C with maximum densification at 550°C. In the present work, a much lower reaction temperature was recorded (460 - 480"C). This can be attributed to the high Al content corresponding to NiAl stoichiometry and the more thermodynamic stability of NiAl compared to NigAl. Slow heating rate (5°C/min) increases the contribution from solid state interdiffusion to intermetalic phase formation at the interface between contacting Ni and Al particles. This inhibits subsequent reaction and increases the temperature required for its initation from 460°C to 545°C for the present sintering conditions. Reactive Hot Isostatic Pressing Because of the densification difficulties encountered during reactive sintering of NiAl in the presence of Y203 or Cr, reactive hot isostatic pressing (RHIP) was applied to get more dense materials. -144- Figure 7 (a, b) shows the as polished structure of NiAI fabricated from 85% elemental (Ni + Al) and 15% prealloyed NiAI with and without Y203 . Density measurements by Archimedes' principle gave 100% of theoretical. The absence of porosity is also shown by the photomicrographs of figure 7. This figure also, shows the uniform distribution of Y203 after RHIP. The etched microstructure (Fig. 7 c,d) shows that the RHIP process caused more refining of the grain size. For RHIPed monolithic NiAI the grain size is about 18 urn comapred to 45 um for reactive sintered material. Considerable grain size refining is observed for RHIPed NiAI-Y203 composite in comparison with the reactive sintered one (Fig. 7d and 3b). In fact, the grain size of RHIPedNiAl-Y203 composite could not be realistically determined because of the very diffuse nature of grain boundaries, particularly for high Y203 content. In general, the microhardness level of the RHIPed NiAl-Y203 composite is higher than that for reactive sintered material (Figs. 2, 8). The microhardness increased also with increasing Y2Og content, however the effect of YgOo is more pronounced for reactive sintered material. Thus, the addition of 15V% YgOo increased the microhardness of reactive sintered NiAI by about 130 VHN, while the same amount of Y203 caused an increase of 80 VHN for the RHIPed material. The microhardness of prealloyed NIAI powder consolidated by conventional HIPing at the smae conditions is less than that of RHIPed NiAI in the absence of Y2Og in both cases (Fig. 8). An example of the microstructure of RHIPed NiAl-38% Cr with and without Y203 is shown in Figure 9. Both Archimedes principle and metallographic evaluation indicate that the density ranges from 98 to 100% of theoretical. The structure is composed mainly of two phases (light and dark phases.). A ten hours heat treatment at 1350°C did not cause any microstructural changes. -145- QMSI Fig.7. Microstructure of reactive HIPed NiAl a, b: unetched c, d: etched a, c: 0% Y203 b, d) 15 V% Y 0 . -146- • *• O< u o><*> >> +J O (II r) <-l O 10 a u o nft 0 o S § I § 8 I cva NUA 4aami>iiiimi3iN m* a O Ot C*» ft 1 O SB >H < II II ••H Z -"im U 3 -147- The microhardness of this two-phase microstructure is given in Fig. 10. The microhardness of the light phase is generally higher than that of the dark phase and increased with Y2Oo addition in a way similar to that of RHIPed NiAl-Y2Oo composite, while that of the dark phase reamined with no noticeable change. It is evident that the two phases both have higher microhardness than RHIPed NiAl-Y203 material indicating the role of Cr which exists in solid solution in the NiAl and as a second phase. Electron microprobe analysis (Table 2) showed that the major constituent of the structure is a NiAl-rich phase which contains about 4,5% Cr. The second phase is a Cr-rich phase alloyed with about 9% Al and 4,5% Ni. High Temperature Sintering In addition to the transient liquid formed during low temperature reactive sintering (700*0 which plays the principle role in reactive synthesis of NiAl and NiAlCr alloys, an attempt was made to investigate the effect of liquid formed at high temperatures in the Ni-Al system. Compacts having 38% Cr with different Al contents were sintered at 1425°C (above the 1350°C eutetic reaction) for 2 hours. Fig. 11 shows the different microstructures, developed by changing the Al content. For 23% Al, the structure is composed of two phases with about 20% porosity. A similar microstructure was obtained by low temperature reactive sintering but with higher porosity (Fig. 5a). Upon increasing the Al content to 24,4% slight shape change of the compact was observed after sintering due to an additional amount of liquid formed at high temperature. The structure in this case is composed of equiaxed two-phase grain structure (Fig. lib) and the density is increased to 95% of theoretical (Fig. 12). Compacts having 27.5% Al suffered slumping and shape loss after sintering due to the excessive amount of liquid formed in this case. The structure is typical for a cast coarse eutectic structure (Fig. lie). -148- m& SOjim Fig. 9. Microstructure of reactive HIPed NiAl-38% Cr a) 0% YjO b( 10 V% Y2°*y {^ ] Fig. 10. Microhardness (VHN) of reactive HIPed MiAl-38% Cr 1) 0% Y40- 2) 3.35 V% Y20 3) 10 V% Y O -149- r-». F-'NHi *V-1 •. ..^^".-vi -..1 -J-..:. y Cffc> 50 Mm Fig. 11. Microstructure of sintered NiAl-38% Cr (100% elemental powders) a) 23 %A1 b) 24.4% Al c) 27.6% Al. 10L 23.0% 24.4% 27.6% Amount of Al(wt%) ' Ughtphate | I Dark phase eel. density Fig. 12. Microhardness (VHN) and density of sintered (100% elemental) NiAl-38% Cr alloys with different Al contents. -150- Electron microprobe analysis of the alloy with 24,4% Al (Fig. lib) indicates that the structure is composed of NiAl matrix with 11,79% Cr and a Cr-rich phase containing 7,6% Al and 8,32% Ni. By comparison with an alloy having 23% Al and fabricated by RHIP (Fig. 9a) there is a difference in the microstructure in addition to more mutual alloying. This can be mainly attributed to the increased amount of liquid phase which enhances interdiffusion of the alloying elements. Microhardness values of these materials are given in Fig. 12. As previously found for materials processed by RHIP, the light phase is harder than the dark one. The lower level of the microhardness of the alloy containing 23% Al compared to similar composition processed by reactive sintering (Fig. 6) may be due to the more coarse structure formed during sintering at high temperature (see Fig. 5a, Ua). Processing of Hypoeutectic NiAl-Cr Material Cold isostatically pressed compacts containing 20 and 25% Cr (100 % elemental powders) were sintered at 1425°C for 2 hours with a heating rate of 15°C/minute. Reactive hipping was also performed for the 20°% Cr composition. Microstructures of processed compacts are shown in Fig. 13. For materials containing 20 and 25% (Fig. 13a C), sintering at 1425°C for 2 hours led to the formating of a more homogeneous microstructure composed mainly of a dark phase surrounding small isolated islands of white phase. For compacts with 20% Cr, a small degree of shape distortion was observed which may be due to the formation of an excessive amount of liquid. Increasing the Cr content to 25% gave better shape control after sintering. For both compositions, the sintered density is about 95% of theoretical; For the reactive hipped 20% Cr alloy the volume fraction of the light phase (CR-rich) is large compared to the -151- sintered 20 or 25% Cr material (Fig. 13). This is supported by the EDX analysis (Table 2) for this alloy which shows that the NiAl phase is only alloyed by 5,5% Cr compared to 9,7r/r Cr for sintered 25% Cr alloy. This can lead to the suggestion that the conditions of sintering without pressure are more favourable for Cr-diffusion in NiAl phase and consequently better homogenity. The low content of Cr in addition to the high sintering temperature can explain the relatively low microhardness level of this group of material (Fig. 14). The slightly higher microhardness of the reactive hipped 20% Cr is in agreement with these considerations. As mentioned above, this material is processed at lower temperature (1200°C) under pressure and contains higher amount of the more hard Cr-rich phase. Samples of the material containing 25% Cr were subjected to reactive sintering at 700°C for 20 minutes with various heating rates ranging from 3 to 20°C/minute. The samples were subsequently containerless hipped (172 MPa, 1250°C, 3 hours). Near full density was obtained for samples reactive sintered with 20°C/minute. The resulting microstructure is given in figure 15. The microhardness is about 500 VHN. Compared to NiAl-38% Cr material processedby sintering at 1415'C of as atomized prealloyed powder25', the reactive sintered NiAl-25% Cr has the same microhardness and is of similar microstructure. The high degree of densification after reactive sintering and the more homogenous structure obtained can be attributed to the more balanced solubility conditions of Cr in NiAl-phase for the hypoeutectic compositions compared to the eutectic alloy processed by the same fabrication route. The mechanical properties of the materials processed in the present work by reactive synthesis are currently being investigated. -152- Fig. 13. Microstructure of consolidated elemental hypoeutectic NiAl-Cr alloys a, c) 20, 25% Cr sintered at 1425*C b) 20% Cr processed by reactive hipping at 1200*C. -153- NIAI-20O NiAI-20Cr NJAK5Cf [ Ughl phase H Dark phase gggg c«L density Fig.14. Microhardness (VHN) and density of hypoeutectic NiAlCr elemental powders 1, 3) sintered at 1425°C 2) reactive HIPed at 1200°C. Fig.15. Microstructure of reactive sintered at 700°C and HIPed NiAl 25% Cr alloy. -154- References 1. G. Sauthoff. Mechanical Properties of lntermetalli.es at High Temperatures; in High-X Temperatures Muminides and Intcnnetallics, C. Liu et al. (eds.) ASMITMS Metals Park (1990) 329. 2. G. Sauthoff. Intermetallic Alloys-overview on New Materials Developments for Structural Applications in west Germany. Z. Metallkoe, 81 (1990) H. 12, P. 855. 3. A. U. Seybolt. Trans. ASM (1966) Vol. 59, P. 860. 4. D.E. Alman and N.S. Stoloff. Powder Fabrication of Monolithic and Composite NiAl. Int. J. of Powder Metallurgy. 27,1 (1991) 29. 5. K. S. Kumar, S.K. Mannan, J.D. Whittenberger, R.K. Viswanadham and L. Christodoulou, Nickel AluminideI Titanium Diboride Composites via XD synthesis, 1989, MML - TR-89-102 (c). 6. J.D. whittenberger, R.K. Viswanadkam, S.K. Mannan and B. Spressler, J. Mat. Sci. 22 (1990) 35. 7. S.C. Iha and R. Ray, J. Mat. Sci. Lett. 7 (1988) 285. 8. J.I. Walters and H.E. Cline, Metall. Trans. 4 (1973) 33. 9. A. Bose, B. Moore, R. M. German and N.S. Stoloff, J. Metals 40, 9 (1988) 14. 10. W.M. Schulson, Inter. J. Powder Met. 23 (1987) 25. 11. A. Bose and R.M. German, Material Sci. and Eng. A 107 (1989). 12. A. Bose and R.M. German, Advanced Materials and process; 3 (1988) 37. 13. D.E. Alman and N.S. Stoloff. Materials Research Society Symposium Proceedings. Vol. 213 (1991) 989. 14. O. Arkens, L. Delaey, J. DeTavernter, B. Huybrechts, L. Buekenhout and J.C. Libouton, Mater. Res. Soc. Symp. Proc. Vol. 133 (1989) 493. 15. G. Sauthoff, Z. Metallkunde, 77; 10 (1986) 654. 16. R.M. German, Liquid Phase Sintering, Plenum; New York, N.Y. (1985). 17. D.M. Sims, A. Bose and R. M. german, Prog. Powder Metall; 43 (1987) 575. 18. R.M. Germna, A. Bose and N.S. Stoloff, Mate. Res. Soc. Symp. Proc. Vol. 133 (1989) 403. -155- 19. C.J. Quinn and D.L. Kohlstedt, J. Mater. Sci., 19 (1984) 1229. 20. Y. Miyamoto, M. Koizumi and 0. Yamada, J.Amer. Ceramic Soc, 67 (1984) 224. 21. W.H. Baek, and R.M. German, Int. J. Powder Metall., 22 (1986) 235. 22. W.H. Baek and R.M. German, Powder Metall. Int., 17 (1985) 273. 23. R.M. German and J.W. Dunlap, Metall. Trans. A., 17 (1986) 205. 24. HE. Cline and J.L. Walter, Trans., 1 (1970) 2907. 25. W. Diehl and M. Dropmann, IAW-KFA, Jiilich Germany, Nov. 1990, Unpublished work. -156- PLG fs-0 15/io First International Spring School & Symposium in Advances in Material Science (SAMS 94) Cairo, 1520 March 1994 MECHANICAL ALLOYING OF NiAl-Cr EUTECTIC ALLOY WITH NON-METALLIC DISPERSOIDS K.E. Mohamed®, D. Stover*, H.P. Buchkremer* and F.H. Hammad* $ Atomic Energy Authority, Cairo, Egypt, & Institute of Applied Material* Research IAW/KFA-Jullich, Germany Abstract Eutectic NiAl-38 wt% Cr alloy powder was mechanically alloyed with non-metallic dispersoids using high energy planetary milling machine. Dispersoid additions of 5 Vol. %Y203; TiB2 or A1203 were used in the present investigation. After 32 hours of milling the dispersoids were sufficiently crushed to fine particles and embedded in the alloy matrix. Alloy powders in the as-received and mechanically alloyed conditions were consolidated by hot isostatic pressing. Substantial microstructural changes accompanied by considerable hardness increase were observed as a result of the mechanical alloying process. Fairly high hardness is retained after 10 hours of exposure at 1300°C (0.90 Tm). 1. Introduction The mechanical alloying process (MA) was developed in the late 1960s^ as a means of combining y precipitation hardening and oxide dispersion strengthening in a Ni-base superalloy, In-853. Since this time, the commercial applications of MA have centered on the development of oxide dispersion strengthened (ODS) Ni-base superalloys. Pe base ODS alloys and dispersion-strengthened Al-allov* -157- Other applications through MA include synthesis of intermetallic compounds, amorphization of alloys and, to a lesser extent, production of alloys from constituents which are normally immiscible in the liquid or in the solid phase. Several reviews of MA have appeared in recent years^2,3,4! in which these developments have been widely covered. The constituents of alloy powder and MAed in a dry high energy ball charge using high energy milling equipment. Interdispersion of the ingredients occurs through repeated welding, fracturing and rewelding of powder particles. These processes lead to progressive refinement of the internal structure of the powder^ '5\ Eventually a balance between the welding and fracturing takes place and a steady state stage is achieved. At this stage the MAed powder is marked by saturation in hardness and constant particle size distribution^, chemical and microstructural homogenity and is heavily cold worked^3,7l The present work describes an attempt to apply the MA process to the NiAl-38 wt % Cr eutectic system. This alloy has a melting point of 1445°C. It can be considered as condidate for high temperature applications due to its microstructural stability in addition to its good oxidation and hot corrosion resistance offered by high Al and Cr concentration. 2. Experimental Procedure Eutectic NiAl-38 wt% Cr alloy powder of particle size less than 45 |im was supplied by H.C. Starck, Berline, Germany. The alloy powder was blended in turbular mixer with 5 vol.% Y203, TiB2 or A1203. The blended powder was loaded under argon into a tool steel milling vial with 10 mm diameter martensitic stainless steel milling balls. The vials were then closed. The ball to powder weight ratio was 11:1. The MA process was carried out by fixing these vials into a high energy planetary milling machine operated at 300 rpm. The processing time ranged from 5 to 32 hours. The MAed -158- powders were uniaxially cold pressed in mild steel capsules, degassed, electron beam welded under vacuum and hot isostatically pressed (HIPed). The HIPing parameters were 120O°C at a pressure of 150 MPa for 3 hours. Micrbstructural investigations using optical metallography and scanning electron microscopy (SEM) as well as microhardness measurements were made for both powders and HIP-consolidated materials. Vacuum annealing treatment for 10 hours was made for the HIP-consolidated materials in a wide temperature range up tol300*C. 3. Results and Discussion 3.1 Processed Powders The as-received eutectic NiAl-Cr powder was high energy milled with different dispersoids for 5, 16 and 32 hours. Pig. (1) show's the morphology of these powders, individual coarse particles with regular shape can be observed for the as-received powder with an average particle size of about 20. urn Fig. (1-a). Agglomeration of powder is seen after 5 hours of milling wih noticeable reduction in size Fig. (1-b). The nearly constant overall particle size with the progress of milling up to 32 hours Fig. (l-c,d) indicates that a steady state stage is achieved. The etched micrograph of Fig. (2) shows the extremely deformed metastable structure of the powder particles after MA for 32 hours. The microhardness measured on the loose powders is given in Fig. (3). The highest hardness is achieved for MAed powders containing TiB2 particles. The hardness of the powder MAed without addition of dispersoids is considerably higher than that of the as received material, however, it is always less than the hardness of the powder MAed with ceramic particles. Clearly the hardness is a measure of the amount of point and line defects introduced in the powder particles during MA as well as the degree of refinement of each dispersoid and its effectiveness in pinning the developed dislocation structures. -159- .!Q. ?s gini m o>) . le) Fig. 1. Morphology of the powder a) as received b, c# d) after MA for 5, 16/ 32 hours, respectively. -160- 3.2 HIP-Consolidated Materials The alloy powder in the as-received and MAed conditions was consolidated by HIPing. The microstructure of the as-received materials shown by SEM micrograph in Fig. (4-b) consists of two phases. Energy dispersive X-ray analysis (EDX) revealed that the light gray phase is a NiAl-base with about 30 wt% Cr and the dark gray phase is a Cr-rich phase having about 51.5 wt% Cr Table (1). It seems that five hours of MA before powder consolidation is insufficient to achieve a good degree of crushing and mixing of the dispersoids in the metallic matrix of the NiAl-38% Cr alloy. In Fig. (4-c) cuboidal TiB2 particles are still easily detected in the microstructure. After MA for 32 hours dispersoids are no longer detected by optical metallography in the HIPed material Fig. (4d). From Fig. (5-b,c) it is important to note that the microstructure after HIPing is essentially the same for materials MAed with or without ceramic dispersoid. This microstructure has certain features which will be analysed below. In the course of MA, the processed powders are more or less contaminated with iron from steel balls and containers. The overall chemical composition (EDX analysis) of the HIPed materials is given in Table (2). Redistribution of elements between the constituent phases takes place resulting in an increase in the volume fraction of the light gray phase and the formation of small volume fraction of a black phase Fig. (5b,c). Compared to the as received HIPed material, the light gray phase is relatively poor and the black one is highly rich in Cr Table (1). A round dark gray phase with an average diameter of about 8-10 um is seen in the micrograph of Fig. (5-a). The EDX analysis of that phase is nearly comparable (Table (1)) with the overall chemical composition of the as received material (Table (2)) with little difference in the order of 2-3%. The microhardness of that phase (550 kg/mm2) is slightly higher than that of the as-received material (510 kg/mm2). Most -161- Table 1. EDX Analysis for the Main Alloying Elements (wt%) of constituent phases as Revealed by SEM Micrographs. 1HIPed as received HIPed MA* 1 Mass Analysed Hi Al cr Ni Al cr tight gray phase 50.51 19.20 29.99 65.47 23.18 10.51 Dark gray phase 35.61 12.29 51.43 45.61 13.36 39.49 Black phase - - - 23.54 10.28 63.75 • Average of 5 different spot analysis Table 2. EDX overall chemical composition of HIPed NiAl-Cr eutectic powder MA for 32 hours with different dispersoids. 1 Amount of Powder dispersoid Chemical composition processing added wt. t Vol.* Iwt.% Ni I Al cr Fo Ti y as 41.72 15.22 42.36 0.65 0.05 O.00 received HA-no dispersoids 0 0 39.69 17.77 41.54 0.88 0.09 0.03 MA with 5 1.8 40.19 16.59 41.47 0.37 0.14 1.24 Y»°3 MA with TiB2 5 2.96 40.71 14.39 41.18 1.06 2.31 0.34 MA with 5 1.63 41.26 15.8 42.09 0.68 0.00 0.17| -162- Ill Fig.2. Etched micrographs of the powder a) as received, b) after MA for 32 hours. 1:00 1 • P ^ 1CC0 I 2 £ E30 17 M "1 1 c 1) a 1 > .* c O u £•:•: 5 ' 01 0 > •- < f C 1 1 1 HI 25 "s* iCG < i-s Fig. 3. Room temperature Microhardness of loose powders. -163- (c) 100C X mmmm fd) jooc r Fig. 4. Microstructure .of HIPed materials a,b) optical and SEM of the as received material, c,d) optical metallography after MA with TiB£for 5 and 32 Hours, respectively. -164- Fig.5. Microstructure of HIPed materials a,b) optical and SEM microstructures for materials MA without dispersoids c) SEM for materials MA with Al^Oj -165- probably this phase represents a small volume fraction of poorly processed particles (or minute areas in particles) remaining after extended MA for 32 hours. White submicron (= 0.33 jxm) particles are seen in the SEM micrographs at 3500 magnification. This phase is seen in all materials MAed for 32 hours even if no dispersoids were added during the MA process Fig. (5-b). The volume fraction of this phase increased when A1203 was addad as a dispersoid Fig. (5-c). EDX analysis revealed that this phase is of high Al content. This phase is believed to be A1203 formed during the MA process due to the presence of an unavoidable amount of 02 in the milling atmosphere although the milling containers were filled and closed under argon. In-situ formed AlnOg has been also observed in some 2 previous investigations^ -°> ® after MA. about 3-7 vol.% A1203 was found to be introduced in the metallic matrix during MA through internal oxidation of the processed alloy depending on the milling conditions and the particular alloy powder composition, Fig. (6) shows the room temperature microhardness of the HIPed materials consolidated from powders MAed for 5 and 32 hours with or without dispersoid additions. Results for the as HIPed an after 10 hours exposure at 1200°C are given in Fig. (6-a,b) respectively. Obviously the hardness in both conditions depends on the processing time and the type of dispersoid added. In general, strengthening in a MAed material containing considerable amount of dispersoid particles which will not be cut by dislocations may be caused by grain boundary strengthening and by the Orowan mechanism of dislocation strengthening^. The increase of strength is proportional to the amount of point and line defects developed during the MA process as well as, the average particlee size and spacing of the dispersoid particles. These particles can serve as pinning points for the non-equilibrium dislocation substructures developed in addition to grain boundary strengthening110^. -166- Evidently, for the present investigation powder processing for 5 hours is insufficient to ahcieve the desirable microstructural changes. Non of the materials MAed for 5 hours has attained remarkable hardness increase over that of the as received material Fig. (6-a). Moreover, the slight hardness increase is nearly completely recovered after 10 hours of exposure at 1200"C Fig. (6-b). This can be attributed to insufficient amount of cold work imparted to the powder particles during MA, as well as poor crushing and mixing of the dispersoid particles. In Fig. (4-c) it was shown that TiB2 particles were still metailographically detected after MA for 5 hours. The effectiveness of ceramic additions becomes more pronounced after MA for 32 hours Fig. (6-a). The highest hardness was obtained for the material containing TiB2 particles and those with addition of A1203 showed the lowest hardnessi ncrease. It should also be noted that considerable hardness increase was obtained for the material processed without addition of dispersoids. As shown above, this material develops a given volume fraction of in-situ formed A1203 particles during the MA process. These secondary dispersoid particles have measurable influence on the hardness of the materials^9K The contribution of such particles to the hardness observed in Fig. (6) remains constant throughout the investigated materials. The changes in hardness measured in this study are due to variation in the primary dispersoid parameters. The materials processed for 32 hours retain fairly high hardness after 10 hours of exposure at 1200°C Fig. (6-b). The different effectiveness of dispersoids depends on their different sensitivity to high temperature, fig. (7) presents the room temperature microhardness after annealing at various temperatures for 10 hours. The microhardness decreases gradually with rising annealing temperature up to 900°C and decreases rapidly above 1000°C. However, at 1300°C (0.91 Tm) the materials MAed with TiB2; A1203 or mA -167- I MA (or 5 hours 8no - D MA (or 32 houri I II •^g 700 E \O l * 600 — M spio s c «= soo TJ _ O Im £.a M 1 CO O cr o O ve d spe r O a CO O I Z'OO '€ >- < o £ < 0 c 1 i 1 1. 1 1 l 1 c l/l 4 < < < < < < Z Z Z 300 < Z Z z * z 1 (a) 700 CM F F I n Ol 600 i/i in ca> •o 500 L. o TJ n O O O CO cro h 600 > O CO ex 1- >• H- z i-2 >• < < 1 1 I 1 1 w«/ i < * .MA - S < < < ispe r < < < 5 •o Z 300 < '•6 z z z z (b) Fig. 6. Room temperature microhardness results for MA materials a) as HIPed condition, b) after annealing for 10 hours at 1200 *C. -168- o As received • MA-no dispersed* A MA-Y203 • MA TiBj O MA AI2O3' 300 200 400 600 6CO 1000 1200 1«X) Anneoling temperature ,°C Fig. 7. Room temperature microhardness after vacuum annealing for 10 hours. -169- without additions of dispersoids still retain about 50% of the hardness achieved before annealing compared to the as recieved material Fig. (7). The gradual decrease of hardness is due to the increased frequency of lattice defect annihilation with increasing temperature'11'. The dispersoid introduced in the metallic matrix during MA inhibit dislocation movement and slow down the recovery processes. Thus, high density, strong dislocation structures created during MA are maintained at high temperatures'12'. The rapid reduction of hardness above 1000°C Fig. (7) should be explained by the decrease of capability of different dispersoids in pinning dislocations and the occurrence of defect annihilation. Although Y2O3 particles are more 8 thermo-dynamically stable than A1203 ' ' and more effective as a hardening dispersoid Fig. (7) the hardness of the MAed material containing Y203 particles is recovered earlier and more rapidly. It was pointed out that the Y2Og phase in the compact alloy matrix containing Al does not exist as Y203 but 8,13 forms with A1203 mixted oxide compounds' '. Formation of the mixed oxide particles starts at the stage of powder cosolidation (T > lOOO'C) and is expected to lead to enlragement of the particles especially after prolonged annealing at high temperatures. This may explain the relatively rapid hardness recovery for the MAed material containing Y203 particles. 4. Conclusions By applying MA process for 32 hours using the present milling condition, ceramic particles were introduced in the matrix of NiAl-38% Cr alloy. Substantial microstructural changes take place after MA. TiB2 particles are more effective as hardening agent in this alloy compared to Y203 or A1203 . In-situ formed A1203 particles have measurable influence on the hardnessof the investigated alloy. Further work is needed to optimize the TiB2 addition. -170- Acknowledgments The present work was performed in the Institute of Applied Materials Research, IAW/KFA Julich, Germany. One of the authors (K.E. Mohamed) would like to acknowledge the financial support of the International Office in KFA during his stay through the bilateral agreement between Egypt and Germany. References /. J. S. Benjamin, Metall. Trans. 1 (1970) p. 2943. 2. R. Sundareson and F.H. Froes, J. Metals, 39 (1987) p.22. 3. P.S. Gilman and J.S. Benjamin, Ann. Rev. Mater. Sci. 13 (1983) p. 279. 4. C.C. Koch, Ann. Rev. Mater. Sci. 19 (1989)p. 121. 5. S.K. Kang and R.C. Benn, Metall. Trans., 18 A (1987) p. 747. 6. J.S. Benjamin and T.E. Volin, Metall. Trans., 5 (1974) p. 1929. 7. J.S. Benjamin, Scientific American, 234, 5 (1976) p. 40. 8. I.S. Poljkin, E.V. Ivanovo and B.P. Matyhin, Structural Applications of Mechanical Alloying, Proceedings of an ASM International Conference, Myrtle Beach, South Carolina USA 27-29 March (1990) p. 131. 9. J.S. Benjammin and M.A. Bomford, Metall. Trans., 5 (1974) p. 615. 10. D.G. Morris and M.a. Morris, Structural Applications of Mechanical Alloying, Proceedings of an ASM International Conference, Myrtle Beach, South Carolina, USA, 27-29 March (1990) p. 265. 11. N. Uenishi and Y. Takeda, Structural Applications of Mechanical Alloying, Proceedings of an ASM International Conference, Myrtle Beach, South Carolina, USA, 27-29 March (1990) p. 41. 12. W.J. Molloy, Advan. Mater. Process. 10 (1990) p. 23. 13. J.D. Whittenberger, Metall. Trans. 15A, 9 (1984) p. 1753. -171- First International Spring School & Symposium in Advances in Material Science (SAMS 94) Cairo, 15-20 March 1994 THE ALUMINIUM-OXYGEN COORDINATION OF ANODICALLY FORMED OXIDE LAYERS ON PURE ALUMINIUM Saleh M. El-Mashri Arab Atomic Energy Agency P.O. Box 402- Tunis 1004, TUNISIA 1. Introduction i The oxide films formed on aluminium by anodic polarization in an electrolyte areUmportant technologically because anodizing treatments u^ed in many engineering applications depend on it. These have a complex microstructure. Those formed on iluminium and aluminium alloys vary from relatively simple films of uniform thickness, when the anodizing treatment fs carried out in neutral electrolytes at low cell voltages,? to very complex porous layers formed in strong acid electrolytes^. In this paper the results of measuring the Al-0 bona lengths in various types of anodic alumina films are reported. These anodic oxide films are prepared in different" electrolytes, using the technique of electron-yield extended X-ray absorption fine structure (EXAFS), in order to determine the molecular structure in terms of the state of coordination of such films. 2. Experimental * The electron-yield measurements from the oxidized surfaces were made over a range of incident X-ray energies extending from 1500 eV (to include the Al K-edge absorption) up to about 2000 eV. This was carried out using the ACO storage ring and double-crystal sort X-ray monochromator operated by Centre National Research Scientifique, Orsay Laboratory, France. Details of experimental arrangement -173- and method of counting the electron-yield are given elsewhere'2'. The oxide films discussed in this paper were prepared by anodic oxidation of pure aluminum, formed under a wide range of anodizing treatment [sodium tartrate, sodium borate, sodium oxalate, chromic acid and phosphoric acid]. More details of the sample preparation and procedure are given elsewhere'1'31 The EXAFS data has been analysed to yield Al-0 bond length by the conventional method of taking the Fourier transform of the EXAFS function using a crystalline oc-alumina as a reference structure to allow for the phase shift. The uncertainty in the values of the bond length has been taken into consideration, and is estimated to be ± 0.0025 nm. This has been discussed in detail^. 3. Results and Discussion Fig. (1) shows the weighted EXAFS oscillations function, X(K)k3, derived from the electron-yield spectra of the oxide formed in neutralized sodium oxalate as discussed in a previous paper^l. The fourier transform of this is given in Fig. (2). The window for the Fourier transform was placed between 2.9 A0"1 and 8.5 A01. The large peak in the transform at R = 0.143 nm is attributed to the Al-0 nearest-neighbour bond length. This needs to be corrected for the phase shift, which was measured from the EXAFS data using standard known bond length of a-alumina'2^. The amount of the phase shift found to be 0.042 nm. Therefore the average Al-0 nearest neighbour bond length is 0.185 nm. It has been found from the EXAFS analysis of the various anodic films, formed in neutralised sodium tartrate and sodium borate electrolytes that they have an average Al-0 bond length of 0.190 nm, while the porous oxide films formed in strong electrolytes; chromic acid and phosphoric acid have on average Al-0 bond lengths of 0.183 nm and 0.180 nm respectively. These bond lengths achieved from the EXAFS analysis for all oxides are summarized in Table (1). The very uniform films formed in sodium tartrate -174- electroyte can be detached from the aluminium substrate by dissolving the substrate in mercuric chloride solution. These have been examined in a TEM fitted with an EELS spectrometer to derive the extended energy loss fine structure (EXELFS) measurements for the oxygen K-edge at 535 eV. The analysis of the EXELFS spectra yielded Al-0 bond length of 0.189 nm for the oxide film formed in sodium tartrate electrolyte, confirming the measurements of 0.190 nm for the Al-0 bond length obtained from the aluminium K-edge of the electron yield EXAFS reported above. More detailed information^ are reported elsewhere. There is clearly a significant difference in the derived Al-0 bond lengths of the different oxides even when the uncertainty in the measured values of the Al-0 bond length is taken into account. This reflects the variation of the molecular structure when different methods of preparation are used. The bond lengths listed in Table (1) show that the oxide films grown in neutralised electrolytes have a significantly greater Al-0 nearest neighbour bond length than those formed in strong acid electrolytes. These bond lengths can be used to deduce information at the molecular level about the state of coordination of the Al3+ ions. In this study where it is intended to develop the relationship between the measured bond lengths and the coordination number for a series of crystalline and amorphous aluminium oxides, a theoretical approach similar to that first adopted by Norman et al.^1. and later by El-Mashri et al.^, is used. For solids with predominantly ionic bonding the crystal potential, according to Pauling,H>1 can be expressed as : e = -Ae2R1 + Be2R"n (1) where A is the Madelung constant, e is the electronic charge, B is repulsive coefficient, R is the nearest neighbour interionic distance and n is the Born exponent. If we consider the crystal at absolute zero, the equilibrium condtion requires e to be a minimum. Differentiation of equation 1, to find the condition for a -175- minimum, d e/ dR = 0, results in . 2 2 2 1 -Ae (- l)Rfl + Be (-n) R*' = 0 which leads to Bn^nl Rn*»-{?= -T-} (2) For similar structures the Madelung constant (A) changes very little^, whilst the repulsive coefficient (B) is proportional to coordination number (N), hence equation 2, may be written as : *-{Sf For the two structures x and y, the ratio of interionic distances is therefore given by Rj lNyJ Since ot-alumina was used as a model compound to calculate the phase in the EXAFS function of the various oxides, because its structure is well known and has approximately the same chemical composition as the amorphous solids, it can again be used here to obtain the coordination number for AF+ in the amorphous oxide using equation 4, taking the interionic bond length for a-alumina to be R = 0.1915 nm, and NY = 6 for reference (All Al3+ ions are octahedrally coordination in a-alumina). In this study the Born exponent (n) has been chosen to take two values; n = 9, according to the analysis used by Norman et al^, and n = 8 which could be considered to be -176- 7 more appropriate for the electron configuration in Al203* l It is now possible to obtain a simple relation between the measured bond length and the corresponding coordination number. Since the EXAFS for an amorphous oxide gives an average bond length, the coordination number which is deduced in this way is an average representation of the structure. If it is assumed that the amorphous structure is simply a bimodal distribution of octahedral [AlOe] and tetrahedral [A104] bonding, as reported by Oka et a\.^ and Norman et al,^51, it is possible to derive the fractions, a and b, of these two coordination states from the average coordination number: N* = 6a + 4b (5) Using equations 4 and 5, a simple relation between the measured bond length of a particular oxide and the percentages of the octahedral and tetrahedral bonding can be derived, as well as the average nearest neighbour coordination. A graphical means of doing this directly from the measured bond length, based on equations 4 and 5 is shown in Fig. 3(a and b). As can be seen in this figure, the two values of the exponent used (n = 9 and n = 8) give quite close results for the coordination state; the vlaues diverge more strongly at shorter bond length. Because of the uncertainty concerning the appropriate value for the Born exponent (n = 8 or 9), the coordination state deduced from Fig. (3 a and b) is taken as the average value of the two curves. In addition to this systematic error there is an uncertainty in the coordination state arising from the error in determining the bond length from the EXAFS. Using this approach for the uniform oxide layers formed on aluminium in sodium tartrate and sodium borate electrolytes, where the EXAFS results give an Al-0 bond length of 0.190 ± 0.0025 nm, it is found that about 90% of the aluminium ions are in octahedral sites and only about 10% are in tetrahedral sites with an uncertainty of ± 8%. From Fig. (3-a) it follows that the average oxygen coordination number for aluminium is 5.66 ± 0.40. The film formed in oxalic acid electrolyte, having an Al-0 bond length of 0.185 ± 0.0025 nm appears to have about 30 + 8% of the aluminium -177- ions in octahedral sites and about 70 ± 8% in tetrahedral sites, while the average coordination number is 4.63 ± 0.40. The aluminium oxide formed in chromic acid electrolyte having an Al-0 bond length of 0.1825 ± 0.0025 nm, appears to have most of the aluminium in tetrahedrally bonded sites and only a small fraction in octahedral sites, with a ratio of 6-fold to 4-fold coordination in the range of (10/90) with an uncertainty of ± 8%. The average oxygen coordination number for aluminium in this type of oxide is 4.10 ± 0.40. As the Al-0 bond length decreases the average coordination number also decreases, which means that more of the Al3+ ions are tetrahedrally bonded. In the extreme case of the oxide formed in phosphoric acid electrolyte, where the measured bond length is 0.180 ± 0.0025 nm, all the Al3+ are tetrahedrally coordinated, or even in a state of lower coordination. There is clearly a wide variation in coordination in amorphous alumina, reflecting the variation of molecular structure when different methods of preparation are used. The average coordination number, and the ratio of [AlOg] / [A104] for the various oxides are summarized in Table (1). Table 1 Electrolyte Al-Onm Calculated oxygen Calculated Ratio used neighbor to Al [AlOg]/ [A104] Sodium Tartrate 0.190 ±0.0025 5.66 ± 0.40 [90/10] ±8 Sodium Borate 0.190 ±0.0025 5.66 ± 0.40 [90/10] ± 8 Sodium Oxalate 0.185 ±0.0025 4.63 ±0.40 [30/70] ± 8 Chromic Acid 0.1825 ± 0.0025 4.10 ± 0.40 [10/90] ± 8 Phosphovic acid 0.180 ± 0.0025 4.00 ±0.40 110/100] ± 8 -178- Fig. i ? Fig.2 0-5 10 15 20 2-5 30 DISTANCE (A) Fig.l The EXAFS function X(k) weighted by K^ for the 50 nm oxalate formed oxide film on pure aluminium. The crosses represent individual data points: the continuous line is the cubic spline fit used for the Fourier transform. Fjg.2 The Fourier transform of the EXAFS function snown in tig. I. -179- 4. Conclusions The structure of amorphous aluminium oxide might consist of a mixture of [AlOg] and [A104] coordination with different proportions depending on the preparation method and on the electrolyte used. This ratio also depends on the purity of the aluminum substrate and the purity of electrolyte, as well as the incorporation of electrolyte anions which can affect structure. For example, the incorporation of an element of higher valency such as phosphorus or chromium should change the bonding in the oxide and therefore lead to change of the ratio of AlOg to A104. Acknowledgments The author likes to thank Professor A.J. Forty, Vice Chancellor, Stirling University and the Staff of the physics Department at Warwick University for their useful and excellent discussions. References 1. S.M. El-Mashri, 1985a, Scanning Electron Microscopy I, 157-163. 2. S.M. El-Mashri, KG. Jones and A.J. Forty, 1983, Phil. Mag. A., 48, 6654-683. 3. S.M. El-Mashri, 1985b, Ph.D. Thesis, Warwick University. 4. A.J. Bourdillon, S.M. El-Mashri and A.J. Forty, 1984, Phil. Mag. A. 49, 341-352. 5. D. Norman, S. Brennen, R. Jaeger and J. Stohr, 1981, Surface Science, 105, L297-306. 6. L. Pauling, 1960, The Nature of the Chemical Bond, (Cornell University Press: Ithaca). 7. A.J. Dekker, 1971, Solid State Physics, 121 (MacMillan Press). 8 Y. Oka, TTakahasi, K. Okada and S. Iwai, 1979, J. Non-Crystalline Solids, 30,349-357. -180- First International Spring School & Symposium in Advances in Material Science (SAMS 94) Cairo, 15-20 March 1994 BEHAVIOR OF Ni- AND Fe- BASE ALLOYS IN ARSENIC VAPOR ATMOSPHERE Y. Afifi*, I. Abdelrazek®, S. Agamy*, and S. Krawczynski* $ Atomic Energy Authority, Cairo, Egypt. $ Nuclear Engineering Department, Alex. University, Alex., Egypt * KFA, Juelich, Germany. Abstract Nickel and iron-base alloys are extensively used in industrial applications because they form chromia scales which are highly corrosion resistant. Some Ni-base alloys also contain aluminum. Aluminium oxidizes at oxygen partial pressure lower than chromium. Therefore, internal oxidation of aluminum in the alloy occurs to a small extent below the surface of the alloy. The alumina is present to stabilize the ordered strengthing phase (NiAl, Ti) and are useful in that they assist the establishment of the protective surface layer of chromia. Addition of cobalt in nickel-base alloys is believed to be the improvements of mechanical properties of the alloys at high temperature. In case of sulphidation of Ni-base alloys, which contain Co as a constitutent, it was concluded that cobalt can alleviate the basic fusion process through the inhibitation of diffusion of sulphur in the alloy. Oxidation does not present a broblem for Ni-base alloys providied that operating temperatures are maintained at appropriate levels. The phosphidation of Ni-Cr alloys resulted in rapidly phosphidized Ni in preference to chromium, while chromium remained unreactive to phosphorous and was concentrated at the alloy side, when the concentration of chromium attained a certain value in the alloy side of the -181- phosphide/alloy interface, chromium was also phosphized, resulting in the formation of (Ni, Cr)-P layers. The additives in the Ni-base alloys, for example Ti/Al ratio and Mo content, play an influenced role in hot corrosion. The Ni contect in Fe-base alloy has a great effect on the corrosion rate. Fe-Cr-Ni-Al alloys are modified at temperatures up to 1300°C. The aim of this work is to study the high temperature corrosion of some Ni- and Fe- base alloys in arsenic atmosphere. This work was done to define which alloys could be used as candidate material in constructing an equipment for GaAs single crystal pulling using Czochrolski method. Hastelloy X, Inconel 617, Inconel 600, and MA 754 were tested as Ni-base alloys. The corrosion temperature was 650°C or 700°C. It was found that increasing the Ni content in Ni-base alloys resulted in increasing the wt. gain%. Ni is the main constituent element in the scale. Cr does not form a protective layer as that for oxidation. Increasing Fe concentration in the subscale is accompanied by decreasing the As concentration and vice versa. Fe-base alloys were tested in As atmosphere for 10, 20, and 30 days at 700°C. The formed scale on Incoloy 800-H was changed into powder as a result of corrosion in As atmosphere. The first subscale of the corroded layer was a porous type. When the corrosion time was increased into 20 days this porous subscale disappeared but the formed layer as a result of corrosion for 30 days as 2 cycles, the first for 10 days and the second for 20 days, contained the porous subscale. The corrosion of MA 956 and Fecraloy had the same corrosion mechanism. 1. Introduction Nickel-base alloys are extensively used in industrial applications because they form chromia scale which are corrosion resistant^1'2'. Nickel-base alloys also contain -182- aluminum. Aluminum oxidizes to alumina at lower oxygen partial pressure than chromium. Therefore, internal oxidation of aluminum in the alloy occurs to a small extent below the surface of the component. The additives of Ni base alloys play an important pole in corrosion^>4,5,6,7] Iron-base alloys are widely used in chemical and petrochemical industries. The corrosion resistance of Fp-base alloys depends essentially on the formation of Cr203 scale16"10]. 2. Materials and Damage Simulation High temperature corrosion in arsenic vapor atmosphere was carried out in a glove box MB-150-GB, mBraun-Germany. Tb i evacuated and sealed tube method was employed for studying the high temperature corrosion of different types of alloys in an atmosphere of arsenic vapor. The samples and the solid arsenic were placed at both sides of the quartz tube, which was sealed under vacuum, according to the following steps. The quartz tube, the samples holder and arsenic boat were cleaned in acetone bath followed by ethyl alcohol and dried using nitrogen gas. The samples were arranged in the sample holder and inserted into the quartz tube followed by the arsenic boat. The quartz tube was sealed with the quartz end plug and then loaded to the glove box. The quartz tube was purged and evacuated many times using the inert gas of the glove box and then was evacuated to 10"5 mbar. During the evacuation process the tube was heated up to 200° C and held for 2 hours using wrapped heater. After cooling the quartz tube was kept under vacuum, 10'5 mbar, for 10 hours. The evacuated quartz tube sealed in the distance between the quartz tube and the end plug providing that its length must not exceed 30 cm, which is the length of heating zone, to avoid the condensation of arsenic vapor during the test. The sealed -183- quartz tube was recharged again into the glove box and then was inserted into the graphite tube. The graphite tube was used to improve the heat distribution at each zone. The graphite tube is inserted into a stainless steel tube, used as a safety tube to contain arsenic vapor in the case of accident, which is blind from one side and the other side had a flange to be connected to the vacuum pump through a valve. The tube was installed in the furnace and then evacuated. A 3 zone furnace, (Gero, Germany), with 3 separate controller units, (Eurotherm Controller 810- Germany) was used to heat up the tube. The samples were tested at a temperature of 700°C, zone 3, for 10 days and some of them were also tested at a temperature of 650"C for 10 days to define the effect of corrosion temperature on the corrosion rate. Some samples were also tested for 20 days at 700°C to define the effect of corrosion time. The arsenic, zone 1, was heated to 617°C for all the experiments, the arsenic vapor pressure was 1 atm at this temperature. Zone 2 heated up to a temperature which was maintained, zone 3, at 700°C or 650°C according to the desired corrosion temperature and zone 1 at 617°C. The temperature at the middle point of each zone was measured using Ni-Ni Cr thermocouple fixed in the graphite tube. Figure (1) shows the experimental set-up for corrosion experiments. The corroded samples were weighed to determine the weight change and then were analyzed metallographically. A scanning electron microscope attached with Energy Dispersive x-ray Diffraction Analyzer, (EDAX), (Philips-4200- Holland) was used to determine the overall concentration profile of each constituent of the alloy. To determine the overall concentration profile of a corroded specimen depletion zone, x-rays were counted whilst the electron beam moves along straight lines parallel to the corroded surface. This averages out any inhomogeneities in -184- the microstructure, and gives a weighted average concentration for a specific distance from the surface. The investigated samples were, Hastelloy X, Inconel 617, Inconel 600 and MA 754, as Ni-base alloys, Incoloy 800H, Fecraloy, and MA 956 as Fe-base alloys. Table (1) gives the chemical composition of these alloys. 3. Results and Discussion 3.1 Results 3.1.1 Nickel-base alloys 3.1.1.1 Hastelloy X It was obvious that the scale was not continuous and spallation of the scale was clear. From these observations it could be stated that the scale is loosely adherent. Cross section examination of Hastelloy x samples corroded in arsenic atmosphere for 10 days at temperature of 650°C and 700°C revealed that the scale structure is a tetra type for both temperatures, Fig. (2-a,b) and the thickness of the corroded layer increases with raising the corrosion temperature from 650°C, the scale thickness is 140 (im, to 700°C where the scale thickness is 190 ^im. Fig. (3) and (4) show the EDAX analysis of the formed scale at 650°C and 700°C, respectively. Ni concentration decreased from 43 wt.% in the first subscale to 38 wt.% in the forth one, i.e., its concentration decreased toward the alloy matrix. Ni concentration in the outermost subscale is 0.9 of that in the alloy matrix for samples corroded at 650"C while it is 0.79 for samples corroded at 700°C. These observations indicated that Ni diffused to the outer surfaces of the samples which are in direct contact with arsenic. Chromium was detected in all the subscales but its concentration increased with increasing the scale thickness for samples corroded at 650°C but only in the first three subscales for samples corroded at 700°C. The maximum concentration of Cr was 1.3 of that in the alloy matrix for samples corroded at 650°C while it was 1.4 for 700°C. -185- Table (1): The chemical composition of the investigated alloys. c Ml •t Cr 11 "Co Tl »1 r< No OtlMOT Alloy \ htulter I t.l ••» o.s ».M 4l.f 1.S0 - - 11 7.0 - TMCPMI «17 •.« - - 13.41 »J.» u.J - l.M - *.« - 1 U (00 *.» X.t 0.1 17.1$ 74.• - - - 7 - - K» 7M - - - !».«• 7».» - • -» 0.1 - - ••"A i*ooi>r ••• i •-•• 0.1 •.» ai.N M.» - o.« a.4 44 - - M *M - - - M.M - - o.s 4.S 73 - ..STO, trzruM .Nt .11 i.a ts.t o.»» O.Q> - 4.74 7f ... °'V, Figure(l): Experimental set-up for corrosion experiments. -186- x 75 (B) x7S Figure (2): The optical micrograph of Hastelloy-X, corroded in arsenic atmosphere for 10 days at: a) 650 C, b) 700 C. -187- I wbtcak: I I I i . HI -J— (21—4- (II •j^frA CO Arstnic so *—* Nicktl a a Chromium ."—• Iron to $301 I *~20 -JT' 10 i.—« !L ^Mi • mi|.» •fcT~I»"T^ I • I i I _1_ _L 0 20 iO SO SO MO 120 HO ISO ISO Oittanct from corrodtd sompk turlac* (jjmj Fir.URE(i): Elemtnlt conctntration in Iht tcalt lormtd on Hasttltoy-X corroded • in arttnk atmotphtit (orttda/s at 6S0'C. 70 U (I) .12). -(31- +-H SO A. —o Arsenic SO A- -» Nkktt O- ••*> Chromium ffc- 10 •* Iron 20 L_. 10 i J—i L_J I i I i L 20 tO SO 10 100 120 UO ISO ISO 200 Dillanet from corrodtd tampit turloctfum) FIGURE (H I: Eltmtntt concentration in tht tcalt lormtd on Hosttlloy-X corrodtd in arstnic atmosphtn for 10 day at 7001 -188- 70 . subicvli - . at ——4—//>- -I3)~-1 SO 50 •—• Artin,c »—i Mc*c/ iO o-• o Chromium -—» Cobalt I jo 20 | 10 _1_ _l_ 20 10 60 60 I0O UO UO 160 HO Distanc* from corrodtd san\tlt surface I//mt FIGURE! *): Eltmtnls distribution in tht scale, fofmtcl on Inconll 6)7 corrodtd in arunic atmosphtrt lor lOUojrs at 650'C. 70 '*-m-\-J2t •{•• Wr4>~w''f—f5JH 60 o—o Arsenic 50 If »_4 N«c 201 10 ^--~ L J, _l_ _1_ "0 20 «7 «0 *ff W0 UO UO /60 Ohtonct tram corrodtd sample- surlocttum) FIGUREIO: Eltmtnts conctntrcthn in tht seal* tormtd pn Ineonel 617 corroded in arstnic atmoxphtrt tor to days at 700 C. -189- Chromium was concentrated in the forth subscale for 650°C and in the third one for 700°C. Iron was peacked in the second subscale for both corrosion temperatures, 650°C and 700°C. Iron concentration was 1.3 of that in the alloy matrix for samples corroded at 650*C and 1.17 for 700'C. The scale was separated from the alloy matrix by a thin layer which contained elements of the alloy but in different concentration than that of the alloy matrix and did not contain arsenic. 3J, 1.2 Inconel 617 Fig. (5) and (6) show the element concentration in the formed scale for Inconel 617 corroded in arsenic atmosphere for 10 days at 65 3.1.1.3 Inconel 600 The formed layer on Inconel 600 corroded in arsenic atmosphere for 10 days at 650*C is a tri-type scale structure while it is a tetra-type scale structure for corrosion temperature of 700°C, Fig. (7-a,b). The elements profile through the scale is shown in Fig. (8) and (9) for 650°C and 700°C, respectively, Fig. (8) shows that both As and Ni concentration decreased through the scale. -190- x OS Figure (7): The optical micrograph of Inconel 600 corroded in arsenic atmosphere for 10 days at: a) 650 C, b) 700 C. - 191 - » submit ~<1)"\r-(>i ^ <*) <\ 60 SO o—o Arsenic «— > NKM to L o-o Chromium n •--» /ron &20 i i 20 fll I iTl I i_J L. X J i L '0 SO «M ISO 200 2S0 100 3S0 iOO Oislanct Irom corrodtd samplt surfocetjim) FIGURE I t):Ettmtnts distribution in tht seal* lormtd on Ineonell 600 corroded in orstik otmosphtrt tor todays atiSO'C ?0I ~ subscatt- t~m—+> 121—4.— 13). -lt)*\ 60 •—o Arsenic SO t—.Hicktl a—o Chromium I, _» • - a Iron 40 1 ; JO 20 f 10 I I • t""l I I _L _L I _L 0 SO 100 ISO 200 2S0 300 ISO tOO Oitlanct from corrodtd samplt surtoctQim) FIGURE(ri.Eltmtnts conctntratien in tht scolt lormtd on lncon*U600 corrodtd in orstnic atmosphtrt lor 10 days at 700 C. -192- Iron was detected in the second and third subscales and its concentration increased from 2.5 to 7 wt.%. Its maximum concentration, in the third subscale, was equal to its value in the alloy matrix. Chromium was detected in the thi/d subscale and its concentration was 1.25 of that in the alloy matrix. Fig. (9) shows that Ni concentration decreased through the first three subscales and Fe was not decreased in the third subscale than that in the second subscale. But in the forth subscale Fe and Ni concentrations were higher than those in the first subscale, the gas/scale interface. Contrary to samples corroded at 650°C, chromium was detected in the last three subscales and its concentration increased with increasing the scale thickness. Fig. (7-b) shows that the morphology of the forth subscale was a columnar type. 3.1.1.4 MA 754 Fig. (10) shows the cross sectional examination of MA 754 corroded at 700°C for 10 days in arsenic atmosphere. Fig. (10-a) revealed that the formed layer is a tetra-type scale structure. Precipitation was observed at the outermost part of the scale, the first subscale, through the second and third subscales and also in the forth subscale, Fig. (10-b). The first and second subscales were separated with a continuous black layer and there were branched transgranular cracks originated at this layer to the matrix, Fig. (10-a). Fig. (10-c) shows that the scale is not a continuous type. Fig. (10-d) shows that the corrosion of arsenic gas is carried out in the first subscale through the grain boundaries. The elements distribution in the formed scale is shown in Fig. (11). As shown Ni is depleted from the matrix to the outer layer of the scale. It is also obvious that As concentration was markedly increased from the first subscale to the second subscale and its concentration is approximately constant in the third and forth subscales. -193- <«) K 10OO (B) N 1000 Figure (10):The optical micrograph of MA 754 corroded in arsenic atmosphere for 10 days at 700 C; a) The formed layer is a tetra type; b)precipitates through the scale; c)The scale is not a continuous type; d) Corrosion in the first subscale is carried out through the grain boundaries. -194- subscali Ill +Y^+ 13) >\ •—-• Arstnie t=J »—4 Nickel o—a Chromium -••••-£ J I I L. 0 20 iO to BO ICO 120 1(0 ISO Distance from corroded sample surfaceOim) FIGUREM: Elements distribution in the scale formed on MA 7St corroded in arsenic atmosphere for »days at 700'C- -195- Chromiui 1 was also depleted to the surface but at a lower rate in comparison with Ni. Titanium and aluminum were depleted for the outermost subscale in concentrations of 1.5 and 2.5 wt.%, respectively. Yttrium and yttrium oxide, Y2O3, were not detected in the subscales. 3.1.2 Iron-base alloys 3.1.2.1 Incoloy 800-H During removing the corroded samples from the sample holder it was noticed that the surfaces of the corroded samples were covered with powder and parts of the scale were spalled off. Fig. (12-a) showed that the scale composed of three different subscales. The scale thickness was around 400 \im. It was also shown that the first subscale which had an average thickness of 280 urn, about 0.7 of the scale thickness, was a porous type. Fig. (12-b) shows the scale morphology at the sample corner. The scale at the corner had only two subscales which were the second and third subscale away from the corner, i.e., it did not have the first porous subscale as that away from the corner. The thickness of the second and third subscales in the corner were about 320 fim which is 2.6 times of that away from the corner. It was also shown that the scale near the corner was completely spalled off. Next to the alloy matrix there was a dark line which separated the scale from the alloy matrix and next to it there were precipitates. The formed layer on the corroded samples in arsenic atmosphere for 20 days at 700°C are shown in Fig. (13), Fig. (13-a) shows that the corroded layer is a tetra-type scale structure. The average scale thickness was about 470 urn The porous subscale, the first subscale in Fig. (12-a). disappeared. But the first subscale was not flat and had traces from the porous subscale which formed for samples 196 Figure(12): The optical micrograph of Incoloy 800-H corroded in arsenic atmosphere for 10 days at 700 C; a) The formed scale morphology b) The scale morphology at the corner. -197- Figure(13): The optical micrograph of Incoloy 800-H corroded in arsenic atmosphere for 20 days at 700 C; a) The formed scale is a tetra-type, b) Corrosion carried out through the grains for the first subscale but it was through the grain boundaries for the third one, c and d)The mechanism of spallation and the cracks propagation in the scale. -198- corroded for 10 days. From this observation it could be indicated that the porous subscalc was already formed but it was porous enough to change to powder during handling the samples. It was also shown that the dark line which separated the alloy matrix from the scale for samples corroded for 10 days, Fig. (12-a), was noticed between the third and forth subscales for samples corroded for 20 days. The fdrth subscale was darker than the others which indicated that it had precipitates. Fig. (13-b) shows that corrosion of Inconel 800-H for 20 days is carried out through the grains for the first subscale but it was through the grain boundaries for the third one. Fig. (13-c) and (13-d) shows the mechanism of spallation and the cracks propagation in the scale. Fig. (14) shows the optical micrograph of Inconel 800-H corroded for 30 days as two cycles, the first cycle for 10 days and the second one for 20 days. It was noticed that the formed scale is distorted and the first subscale is a porous type. Fig. (14-b) shows that the corrosion is carried out through the grains in the second subscale and through the grain boundaries next to it, in the third subscale. Corrosion through the grains is hsown again in the forth one. 3.1.2.2 MA 956 Samples of MA 956 alloy were corroded for 10 days. The corroded samples for 10 days were corroded for another 10 and 20 days. Fig. (15) shows that the cross sectional examination of MA 956 samples corroded for 10 days in arsenic atmosphere at 700°C. It is noticed from Fig. (15-a) that the surface of the corroded sample is not more flat but it has a wavy front. The average scale thickness was around 12^m. The formed scale was separated from the alloy matrix with an interface, the black line, which was relatively parallel to the wavy front of the surface of the corroded sample. -199- W) luift Figure (14): The optical micrograph of Incoloy 800-H corroded in arsenic atmosphere for 30 days at 2 cycles, the first cycle for 10 days and the second one for 20 days at 700 C: a) The formed scale is distorted, b) shows that the corrosion is carried out through the grains in the second and forth subscales. -200- M 1230 (B) x 350 Figure (15):The optical micrograph of MA 956 alloy corroded in arsenic atmosphere for 10 days at 700 C; a)The formed scale has a wavy front, b) Precipitates are formed next to the interface inward the alloy matrix. -201- Fig. (15-b) shows that precipitates are formed next to the interface inward the alloy matrix. The formed scale on MA 956 alloy samples corroded in arsenic atmosphere for 20 days ais 2 cycles, each cycle for 10 days, is shown in Fig. (16). The average scale thickness is around 70 |jm. The formed scale was not continuous and the interface between the scale and the alloy matrix became thicker than that formed for samples corroded for only one cycle, Fig. (15) and (16). The formed interface penetrates into the alloy grains forming precipitates, Fig. (16-a). These precipitates are oriented. It is also observed that the formed scale contained cracks which joined the gas environment at the interface between the scale and the alloy matrix. Fig. (16-b) shows that the formed scale is intersected and these intersections are composed of two different structures. One of them is the black area in the middle and the second is the outer shell, the white area. Precipitates were also observed in the second structure, the white area, around the middle black area and near the interface which separated the scale from the alloy matrix. The formed scale for MA 956 samples corroded for 30 days as 2 cycles, the first cycle as for 10 days and the second one for 20 days, is shown in Fig. (17). It was noticed that the formed scale did not contain the intersected layer and the interface between the scale and the alloy matrix, as shown in Fig. (16). From this observation it could be indicated that the intersected layer has completely spalled off after forming. It is also noticed that corroding the sample is carried out through diffusing arsenic gas into the alloy matrix as a result of forming cracks in the formed scale and forming longer precipitates. 3.1.2.3 Fercraloy Fecraloy samples were corroded for 10 days. The corroded samples were corroded for another 10 and 20 days. -202- (A) X 100 (B) X 1O0 figure (l6):The optical micrograph of MA956 alloy corroded in arsenic atmosphere for 20 days as 2 cycles, each cycle for 10 days, at 700 C; a)The formed interface penetrates into the alloy matrix forming precipitates, b) The formed scale is intersected and the intersections are composed of two different structures. -203- 12S Figure(17): The optical micrograph of MA 956 alloy corroded in arsenic atmosphere for 30 days as 2 cycles, the first cycle for 10 days and the second one for 20 days, at 700 C. -204- Pig. (18-al shows the cross sectional examination of Fecraloy samples corroded for 10 days. It shows that the scale thickness is around 6 (am. It is also shown that the formed scale is separated from the alloy matrix by a relatively planar interface. Fig. (18-b) shows the microstructure of Fecraloy sample corroded at 700°C for 20 days as two cycles, each one for 10 days. It was noticed that corroding the samples for another 10 days resulted in intersecting the formed scale and spalling off parts of the scale. It was also observed that the interface between the scale and the alloy matrix became thicker and penetrated into the alloy matrix forming precipitates. It was also noticed that the formed scale above the interface between the scale and the alloy matrix is grown outward into the gas environment and contained precipitates in its middle near the interface. The scale also contained cracks which joined the surface of the sample at the interface. The thickness profile of the scale formed outward the gas environment is proportional to the formed precipitates density in the alloy matrix. The cross sectional examination of Fecraloy sample corroded for 30 days as two cycles, the first one for 10 days and the second for 20 days, is shown in Fig. (32-c). It was found that increasing the corrosion time for the second cycle to 20 days resulted in spallation of the formed scale in the gas environment. It is also noticed that the interface between the scale and the alloy matrix, Fig. (18-b), is not on the outermost surface as expected to be but it is formed cutting the growing precipitates. Its shape is distorted and is thinner than that for samples corroded for 2 cycles, each one for 10 days. 3.2 Discussion It is obvious from comparing the behavior of the studied Ni and Fe-base alloys, the morphology of the scale and the •^calo thickness in different atmospheres with •••-nlt> 20S <« « J73 Figure (18): The optical micrograph of Fecraloy samples corroded in arsenic atmosphere at 700 C for; a) 10 days, b)20 days as 2 cycles, each one for 10 days, c) 30 days as 2 cycles, the first cycle for 10 days and the second one for 20 days. -206- obtained from corrosion in an atmosphere of arsenic vapor presented above, that corrosion mechanism in arsenic gas and the morphology of the scale is completely different than other gases. Arsenic gas is more aggressive than others, oxygen, sulphur and sulphur dioxide and phosphurous gases, The high corrosivity of arsenic could be referred to its different valences, +3, +5, and -5, and most metals form arsenides. The arsenides take the form of MgAs., M5As, M4As, MgAs, M3As2, M2As, M5As3, M3As2, M4As3, M5As4, MAs, MgAs4, M2As3, MAs2 and M3As7. Many of these intermetallic compounds exist over a range of compositions, and nonstoichiometry is obtained^8'. This property of As results in increasing the lattice defects in the alloy which results in proceeding the arsenidation process faster than others. It was also found that the scale is loosely adherent which could be referred to outward metal and the formed arsenides migration through the scale, and development of growth stresses in the scale. The formed scales on the investigated alloys are multilayer type because the constituent elements in the alloys do not have the same chemical potential to As which leads to a selective reaction of As with one of these elements and when the partial pressure of As reaches a certain value in the scale the reaction of a second elements proceeds forming another layer different from the formed one for the first element or another phase in the same layer. Spallation of the formed scale could also be referred to forming nonstoichiometric arsenides which results in increasing the lattice defects in the scale. Heating the investigated alloys at 650°C or 700°C for 10 days results in sintering of the formed arsenides. Because the corrosion is carried out through a selective reaction, the formed arsenides are separated from each other either in layers or in the same layer. Because sintering of these arsenides are expected to be different, the arsenides are sintered at different rates leading to forming microchannels in the scale. These microchannels are filled with As gas which results in -207- separating the sintered parts from each other. These parts are spalled off during cooling down the samples. The formed scale was separated from the alloy matrix with a layer which contained the contributed elements in forming the scale but in different concentration than that in the scale and also the alloy matrix and did not contain arsenic as a constituent. This could be referred to diffusion of the alloy elements outward through the scale to the gas environment. The thickness and the chemical composition of this layer is different from one alloy to another and also for the same alloy when the corrosion temperature is changed because the alloy elements have different diffusion rates and their concentrations are dependent on the order of their reaction with arsenic. 3.2.1 Nickel-base alloys Nickel was the main constituent element of the first subscale. This is attributed to the higher reactivity of Ni with arsenic compared with the other elements, e.g., Cr and Fe. For this reason Ni rapidly reacted with arsenic in preference to the other alloy elements. The Ni-As can take many forms depending on the As to Ni ratio. According to Ni-As phase diagram, nickel arsenidies at temperature of 650°C or 700°C can be found as four intermetallic compounds, (NiAs2 (a and (3 forms). NiAs, Ni^Asg, and Ni5As2 • Two of which are essentially stoichiometric, (Ni^Asg and NiAs2), and the four compounds have different lattice parameters. It is expected that the most arsenic-rich arsenides formed firstly next to the gas/alloy surface interface and the most arsenic deficient arsenides next to the alloy surface. Because of the high diffusivity of arsenic gas into the alloy the Ni arsenides changes from one phase to another starting from the surface in contact with the gas environment and then through the formed scale and could also be found together in the scale. The lattice defects concentration is expected to be high from the beginning of -208- the corrosion process because of the continuous changing of Ni-As system from one phase to another depending on the abundance of As in the scale. Not only Ni-As have more than one phase but also the other elements, as will be mentioned later. For this reason the lattice defects concentration is expected to be higher in case of corrosion in arsenic gas than other gases. 3.2.1.1 Chromium rffect It was clear from the results that Cr did not form a protective scale as that in the other system. This phenomena could be explained as follow: Cr is a lower reactive element, with respect to nickel and to arsenic. For this reason Cr remained ur.reactive to arsenic till its concentration attained a certain value in the alloy side of the arsenic/alloy matrix interface, then it started to form Cr-arsenides. The phase could be expressed as (Cr, Ni) As layer. For this clarification, Cr was not detected in the same subscales for the tested alloys. Cr was peaked into a definite subscale which may be referred to the high lattice defects concentration in this subscale which resulted in diffusion of other elements, such as Cr, to this subscale. This phenomena could also be explained as: this subscale may have larger voids or micro channels concentration than the others which may be filled with arsenic vapor resulting in increasing the arsenic vapor pressure in this subscale than the preceding one. Increasing the vapor pressure results in depleting the less reactive element while the higher reactivity elements had already reacted from the normal diffusive arsenic. 3.2.1.2 Iron effect For Hastelloy X, Figs. (3 and 4), and Inconel 600, Figs. (8) and (9), which contained Fe as a constituent element of the alloys, it was found that decreasing the Ni concentration was accompanied by increasing the Fe concentration and -209- vice versa. It is postulated that Fe decreased the cation vacancy concentration in the arsenide lattice. This is similar to the results obtained for the effect of iron addition in Cr-Ni system corroded in phosphorous vapor atmosphere^. The Fe might be found in different phases as As + As2Fe, AsFe, AsFe/AsFe2, AsFe2 + a-Fe depending on the As and Fe concentration according to the As-Fe phase diagram^1Q1. Fe reacts with As according to the following equations^11^. 2As (g) + Fe (s) > FeAs2 (s) (1) As(g) + Fe(s) >FeAs (s) (2) As (g) + 2 Fe (s) > Fe2As (s) (3) From these equations and from Figs. (3, 4, 8 and 9) it could be postulated that the Fe arsenides in the scale are in the following order: FeAs2 , FeAs and Fe2As from the gas/alloy interface to the scale/alloy interface. Increasing the corrosion time would result in transforming Fe-arsenides from the most arsenic-rich arsenides to the most arsenic-deficient arsenide. This transformation would result in increasing the lattice defects concentration which was followed by diffusing another element to the subscale. Therefore, the subscale in this case was composed of different alloy elements and the reaction was not a selective reaction as in the case of oxidation. 3.2.1.3 Cobalt effect Although Co is more reactive to As than Fe and it was expected that Co reacted before Fe, cobalt was not detected in the scale for Hastelloy X. This might be referred to the low concentration of Co in the alloy matrix, while for Inconel 617 which contained Co in concentration of 12.5 wt.% Co was detected in the scale. 3.2.1.4 Effect of corrosion temperature Raising the corrosion temperature from 650 to 700" C -210- resulted in increasing the corroded thickness and also wt. gain %. This result is expected because arsenic vapor pressure is very sensitive to temperature. At 650°C the vapor pressure in 2 bar while it is 5 bar at 700°C'12'. Increasing the vapor pressure would result in promoting the reaction rate and also the corrosion rate. 3.2.2 Iron-Base Alloys 3.2.2.1 Incoloy 800-H It was noticed that the surfaces of the corroded samples of Incoloy 800-H alloy were covered with powder which could be referred to the formation of a porous subscale. The optical micrograph investigations showed that the first subscale for samples corroded for 10 days and 30 days were a porous type while it disappeared for 20 days. This could be explained as; in the beginning of the corrosion Ni reacts with As and the products change from one form to another, as discussed in section 3.1, resulting in increasing the lattice defects concentration in the scale. Iron is expected to be the second alloy constituent element reacted with As next to Ni. From the EDAX analysis of Ni-base alloys mentioned in the above section it was concluded that increasing the Ni concentration in a subscale is accompanied by decreasing the Fe concentration in the same subscale. This phenomena referred to decreasing the cation concentration in the formed arsenides. Iron also changes from one phase to another and because its concentration is higher than that in Hastelloy X and Inconel 600, Incoloy 800-H alloy has Fe of 44 wt.%, the reaction resulted in increasing the lattic defects in the scale than that for Hastelloy X and Inconel 600. For these reasons, the first subscale is a porous type, increasing the corrosion time for 20 days resulted in the disappearance of this porous subscale which could be explained as: increasing the corrosion time resulted in increasing the thickness of this porous subscale where it was 70% of the scale thickness [or samples corroded for 10 days. Increasing the pores and voids -211- concentration, as a result of increasing the scale thickness, results in changing the porous subscale into powder easier than that for corrosion time of 10 days. Corroding the samples which were corroded for 20 days, for another 10 days resulted in appearing this porous subscale again which could be referred to depleting the JNi.-and-Ee elements from the alloy matrix and repeating the mechanism discussed above. This porous subscale also changed partially into powder which explains the distorted shape of the surface as shown in Fig. (14). 3.2.2.2 MA 956 The corroded sample is not flat but it has a wavy front which could be referred to the outward diffusion of the main constituent elements, Fe and Cr, into the gas environment and different corrosion and diffusion rate of these elements. The formed scale was separated from the alloy matrix with an interface, the black line, which could be referred to depleting one or more of the alloy elements, it is thought to be Al and Ti, and forming arsenides with lower diffusion rate than that of the main constituent elements in the alloy, Fe and Cr. This interface is blocking the alloy matrix and decreasing the diffusion rate of the matrix elements. For this reason the scale has precipitates next to the interface inward the alloy matrix. Corroding the samples for another cycle, 10 days, resulted in increasing the interface thickness which may be referred to depleting more Al and Ti from the alloy matrix. It was noticed also that the precipitates were oriented because of the lower diflusion rate of As into the alloy matrix and carrying on the corrosion through the grains. It was also noticed that the formed scale is intersected and these intersections looked like irregular oval shape and were composed of two different structures. This could be explained as : the interface is weakened in one place which resulted in diffusing As gas into the alloy matrix and - 212 - reacting with the depleted Fe and Cr from the matrix. As the concentration of As gas in the alloy matrix is decreased, the reaction between the As and the less elements concentration in the alloy matrix, Al and Ti, took place resulting in forming another interface. For this explanation the irregular oval shape parts of the scale had a black part in the middle which was a part of the interface and the white area of it also had precipitates which are thought to be Cr arsenides but the main product in the whigh area are thought to be As-Fe products. The formed scale had cracks in it because of changing Fe arsenides from one phase to another as discussed above. Increasing the corrosion time for the second cycle to 20 days resulted in the disappearance of the irregular oval shape parts of the scale because of their complete spallation. This is because of the high surface tension of these sintered parts. 3.2.2.3 Fecraloy The morphology of the scale formed on the corroded Fecraloy samples were roughly the same as that for MA 956 alloy samples because there is not a great difference in their chemical composition. 4. Conclusions 4.1 Ni-base alloys From the obtained results for the corrosion of Ni-base alloys in an atmosphere of arsenic gas it could be concluded that arsenic vapor atmosphere is more aggressive than the other atmospheres. The scale formed as a result of corrosion in an atmosphere of arsenic vapor is a multilayered type. The scale is loosely adherent and parts of it spalled off. Ni is the main constituent in the scale. Cr did not form a protective scale. The scale separated from the alloy matrix with a layer which contains the alloy elements, not all, but in cocentrations different than that of the alloy matrix. Fe -213- decreases the cation vacancy concentration in the arsenides. Increasing Fe concentration in the subscale was accompanied by decreasing arsenic concentration in the same subscale. Co diffusion to the outer subscales depended on its concentration in the alloy matrix. The effect of Y203 on the scale type was not noticed in comparison to its effect in case of oxidation and salt-melt hot corrosion. 4.2 Fe-base alloys Investigating the Incoloy 800-H indicated that the corrosion of Fe-base alloys containing Fe and Ni as main constituent elements in the alloy matrix resulting in forming a multilayer scale. The first subscale of the forming corroded layer is a porous one which is about 70% of the scale thickness. This porous subscale changed into powder as a reuslt of increasing the corrosion time or handling the samples. The edge effect is obvious for this type of alloys. Chromium did not form a protective scale as that in Ni-base alloys. From the results for MA 956 and Fecraloy alloys, it could be concluded that the corrosion of Fe-base alloys which do not contain Ni resulting in a completely different scale types. These scale have a wavy front which changed into irregular oval shape as a result of increasing the corrosion time to 20 days. Increasing the corrosion time to 30 days results in spalling these discrete parts which have the irregular oval shape. Also, for these types of Fe-base alloys, there is an interface between the scale and the alloy matrix. The thickness of this interface increases with increasing the corrosion time. Next to this interface inward the alloy matrix there are oriented-growing precipitates. From these results it could be concluded that the alloys which do not contain Ni as a constituent element as its matrix has Al and Ti as constituent elements could be promising for corrosion in As gas ubt it should be investigated in more details to know the effect of Al, Ti and Cr concentration on the -214- morphology of the scale. This study might be focus on with the effect of these elements on the interface thickness and the type of the precipitates inward the alloy matrix. References 1. A.S. Radiff: Materials science and technology July 1987 vol. 3: pp. 554-561. 2. A.C. Dean : in proceedings of the petten international conferance: JRC petten Establishment. Petten (NH), the netherlands, 15-18 October, 1979: Published by the material society, London, United Kingdom, 1980; pp. 271-296. 3. Guan De-Lin, Lu Zhi-han, Ma Chong-hui, and Xiao Yao-tian: In proceedings of third JIM international symposium, MT. Fuji. Lake Yamanaka, Japan. November, 17-20, 1982 Puplished by the Japan institute of metals Aoba Aramaki Sendai 980, Japan: pp. 401-410. 4. Greig R. Wallwork and John M. Newburn : In proceedings of third JIM international symposium, MT. Fuji, Lake Yamanaka, Japan, November, 17-20, 1982 Puplished by Japan institute of Metals Aoba Aramaki Sendai 980, Japan: pp. 103-114. 5. F. Armant, A. Vejux and G. Beranger: In proceedings of the pMen international conference: JRC petten Establishment, Petten (NH), the nether lands, 15-18 October, 1979: published by JRC petten Establishment, London, United Kingdom, 1980; pp. 423-434. 6. Yoshinori Sasaki : In proceedings of third JIM international symposium, MT. Fuji, Lake Yamanaka, Japan, November, 17-20, 1982 Published by the japan institute of metals Aoba. Aramaki Sendai 980, Japan: pp. 661-668. 7. Dunxu Zou, Zhenlan Chen and Fengjie Gonfj; in Proceedings of third JIM international symposium, "High temperature corrosion of metals and alloys": MT. Fuji, Lake Ya/nanaka, Japan, November, 17-20, 1982 Published by the japan institute of metals Aoba Aramaki Sendai 980, Japan: pp. 377-384. CREVICE CORROSION OF HASTELLOY C-276 AND INCONEL-625 ALLOYS IN CHLORIDE ENVIRONMENTS I: Determination of the Critical Crevice Solution M.MLA Gad, AA. El-Sayed and S.M. EI-Raghy* Dept. of Metallurgy, Nuclear Research Center, Atomic Energy Authority, Cairo, Egypt Cairo University, Faculty of Engineering Abstract The crevice corrosion of HastellOy C-276 and Inconel-625 alloys is investigated. The effect of temperature on both passive current and the critical pH (at which the passive film breaks down) is studied in detail. Potentiodynamic studies have been used to determine the critical crevice solution (CCS) in terms of pH and chloride ion concentration, and passive current. The criteria of both Crolet and Oldfield and Sutton are used for the determination of the critical pH for crevice corrosion. The results show that the critical pH for crevice corrosion of Inconel-625 alloy as defined by Oldfield and sutton seems to be a function of both temperature and chloride ion concentration, whereas the same relationship does not appear to hold between these variables and the critical pH when the Crolet criterion is applied. The critical pH of crevice solution for Inconel-625 may vary between 1.7 and 0.0 depending on environmental condition and criteria applied. However, for the case of Hastelloy C-276 alloy the critical pH for crevice corrosion could not be detected when applying both criteria. The results indicate that Hastelloy C-276 alloy is more resistant to crevice corrosion than Inconel-625 under the same experimental conditions. The experimental results passive current (Ipass) and -217- depassivation pH (dpH) of this investigation could be utilized as inputs to a mathematical model employed for the prediction of incubation period required to reach the critical crevice solution composition, which leads to permanent breakdown of passive film. 1. Introduction Nickel-based alloys have a very wide range of applications. They are used as corrosion resistant materials^ ' , in heating elejpaent manufacture^ ' and also as creep resistant alloys^3'. Other applications of these alloys include construction materials for containers used in radioactive waste disposal^. Nickel-based alloys can form a protective passive film in some environments which leads to improvement of the corrosion reistance. However, these alloys may corrode in environments containing appreciable amounts of chloride or halide ions^, especially if oxidizing species are present. The presence of oxidizing species coupled whith elevated temperatures may destroy the resistance of these alloys in chloride solutions^. The attack may appear in different forms of corrosion e.g. uniform corrosion, localized corrosion (crevice corrosion) and stress corrosion cracking (SCC)t5,6]. Crevice corrosion is a catastrophic form and could be more serious than pitting corrosion. The mechanism of initiation of crevice corrosion of passive metals can be summarized as follows: 1. As corrosion proceeds, the electrolyte in the crevice becomes depleted in oxygen because it can not be replaced sufficiently rapidly by diffusion. 2. This situation creats a galvanic cell where the cathodic reaction involves oxygen reduction outside the crevice (a relatively vast area). The solution in the crevice and an anodic reaction inside the crevice (a very small area) initiating metal oxidation'7'. Migration of CI' into the crevice in order to balance the accumulated +ve charges -218- resulting from hydrolysis of metal chlorides. 3. Finally, the composition of the solution in the crevice reaches a critical value at which the passive film breaks down and rapid corrosion commences. pH values in the range 0-3 and chloride ion concentration up to 12M have been reported for crevices and pits in common SSs17,81. It was Crolet[9] and co-workers who introduced the concept of a depassivation pH or more appropriatly activation pH at which the passive film breaks down and rapid corrosion pccurrs. The depassivation pH is an inherent property which reflects the resistance of metals and alloys to crevice corrosion, the lower the pH at which passivation breaks down, the longer the inititation process will takes place. Determination of activation pH involves the recording of anodic polarization curves at decreasing pH until the current peak exceeds 10uA/cm2. An entire set of measurements can be performed on a single test surface. Each successive test is performed on a surface which has undergone changes in the same way as it would have in real crevice. Oldfield and Sutton^101 have developed a technique based on a mathematical model for the crevice corrosion process. According to this model the process is divided into four stages (deoxygenation, pH fall, breakdown of passive film and propagation). The model requires various parameters as inputs such as alloy composition, crevice geometry, chloride level of the bulk solution and so on. However, the model also requires two inputs which have to be determined electrochemically (critical crevice solution, CCS and passive current, Ipass). The CCS is determined by slightly modifying Crolet's depassivation pH method. In this case the chloride level in the test solution increased as the pH is reduced as what would happen in the real crevices, polarization curves are determined in increasingly aggressive solutions and the CCS is the solution in which anodic active peak height reaches 0.1 A/cnr. The present investigation is devoted to the determination of the electrochemical parameters (Ipass, CCS) which are currently used as inputs for the evaluation of the crevice corrosion by mathematical modelling of two Mo containing nickel-based alloys namely, Hastelloy C-276 and Inconel-625. 2. Experimental Work Cylinderical test specimens (12 mm length and 9 mm diameter) were wet ground on silicon carbide paper up to 600 grade. The chemical composition is given in table 1. The test solution were prepared from doubly distilled water and analytical grade NaCl. Once the appropriate temperature, the solutions were then acidified with HC1 acid to the desired pH (measured at temperature) and deaerated with nitrogen. Studies were undertaken at temperatures of 25, 50 and 100°C in solutions of 0.5, 1 and 4 M NaCl solutions. A conventional multinecked 1 liter glass corrosion cell was used for the electrochemical tests. A saturated Calomel electrode, SCE conected to the glass vessel via a salt bridge was used as the reference electrode. The electrochemical equipment was an EG&G PARC, MODEL 350 A Electrochemical Corrosion System with IR compensation, MODEL 356 EG&G PRINCTON RESEARCH. Potentiodynamic tests were performed using polarization rate of 10 mv/min, and the scans were started at 100 mv below Ecorr and stopped at 0.4 V vs SCE. 3. Results during the incubation stage of crevice corrosion, the underlying metal or alloy remains protected by the existing passive film*1'. In this stage the metal corrods in a slow manner at a rate depending on the properties of the passive film. Once the film is partially damaged active corrosion occurs at a higher rates. The composition of the critical crevice solution (CCS) in terms of pH and chloride ion concentration is determined from a series of potentiodynamic polarization tests in a number of concentrated deaerated -220- solutions as environment inside the crevice. The specific solution pH and chloride level at which an anodic current peack of 10 uA/cm2'10' is measured during potenital scanning is identified as CCS. However, the Crolet criterion^identified the CCS as corresponding to the pH and CI" concentration at which the first peack appears. Figure 1 shows a series of potentiodynamic scans in the temperature range 25-100°C, and pH=0.0 for both alloys. It could be seen that for inconel-625 alloy, there is a clear transition from active to passive state resulting in two current values namely, the peack and the passive currents, while for alloy C-276 no transition is observed. The effect of temperature on both current values is shown in tables (2,3,4) and Figs. (2-10). Figures (2-10) present the results of potentiodynamic scans in 0.5, 1 and 4 M NaCl for both C-276 and Inconel-625 alloys. Applying Oldfield and Sutton^10! criterion on the behaviour of alloy 625 shows that the critical pH ranges between 0.0 and 0.8 depending on temperature. On the other hand, when the Crolet^ concept is applied, the critical pH value may vary between 1 and 1.8 regardless of temperature. However, for C-276 alloy, the active peack current could not be detected from a series of potentiodynamic scans (Fig. 1). It is clear from Figs. (2*10) and Tables (2,3,4) that the increase in both the CI' concentration and temperature lead to a corresponding increase in the values of the currents (Ipeak axA Ipass) and a reduction in the value of pH. 4. Discussion Literature survey indicates that the extensive work conducted on crevice corrosion of SSs and Ni-based alloys reveals the importance of some key parameters to this phenomenon namley, the Ipass and the composition of the critical crevice solution (CCS) at which passivity breaks clown1"'12'. These parameters may be considered as a -221 - rm.is a.una • llnoticlloy C-627 1- 25 C 2 !3 2- SO *C 3- 100 *C -a.MO >&F~- Jo1 id* IS* ij* ids ": iuWcnt ib' io3 'o"4 IOS ios- lo'Nn/cn* jFi#. (X) : Series of potentiodynamic polarization tests in 0.5 MNaCl Solution at pH = 0.0. -222- -'-"--• T—"-- 3B S7. . - u o • .•;." : '' '•.""'; i •: . : :; ;: • ; ii:: ft- ::«;;,!.. I | S^ 3 *«• E- =j w ,s. 1• *•* ' **3 •1 X-r- !iPg 1 O ,'N 1 !tt :••>:' *1 $ '• >M IM si. 1 1 O ! :fc •.<:•:•' —— ^3 «? Ha T* i^^:' MS \ V N \ -V-\ ^> A -- 8 o 8 8 8 8 8 8 § in O m 0 m o u> «» o n w w t- juw / yu ' ssvd | - 223 - Table (1): Chemical composition, wt.% Element Ni Cr Mo Fe C Mn C-276 bal. 15.6 15.6 6.1 0.002 0.48 1-625 bal. 21.73 8.49 4.56 0.03 0.16 Element P S Si Ti W Co C-276 0.005 0.002 0.02 - 3.73 1.22 1-625 0.01 0.001 0.25 0.3 — — Table (2): Peak Current Values of Inconel-625 in 3% NaC! Ipeak , nAJcm2 pH 25'C 50'C 100 V 0.0 1 *10E4 2 *10E5 2 *10E6 0.1 2 * 10E3 1 * 10E4 9 *10E4 0.2 5 *10E3 3 *10E3 7*10E4 0.3 ; 5 *10E3 2 *10E4 8 *10E4 0.5 5 *10E3 1.1 * 10E4 2 *10E4 0.7 1 *10E3 8 * 10E3 2 *10E3 -224 - Table (3): Peak Current Values of Inconel-625 in 6% NaCl. Ipeak , nAJctn2 pH 25'C 50 V 100'C 0.0 2 *10E5 3 *10E6 2 *10E7 0.1 9 * 10E3 8 *10E4 7 *10E6 0.2 8 *10E3 3 *10E4 8 *10E5 0.3 6 *10E3 2 * 10E4 2 *10E5 0.5 3 *10E3 3 *10E4 1 *10E4 0.7 2 *10E3 1 *10E4 9 *10E4 Table (4): Peak Current Values of Inconel-625 in 24% NaCl. Ipeak , nAJcm2 pH 25'C 50'C 100'C 0.0 6 * 10E6 1 * 10E6 2 * 10E7 0.1 9 * 10E5 4 * 10E5 5 * 10E6 0.2 9 * 10E4 1 * 10E5 1 * 10E6 0.3 9 * 10E3 4 * 10E4 5 * 10E5 0.5 7 * 10E3 2 * 10E4 9 * 10E4 0.7 4 * 10E3 6 * 10E3 2.1 * 10E4 -225- measure of the alloy resistance to the initiation of crevice corrosion. Recently, mathematical models^13,14,15! were developed where these parameters were fed along with fixed parameters (alloy composition, crevice geometry and bulk solution composition) to predict the time to breakdown of a passive film. The results of the present investigation will be discussed in terms of the value of the passive currents and the depassivation pH (dpH) at which the film breaks down to rank the studied alloys according to their susceptibility to crevice corrosion. Critical Crevice Solution: Critical crevice solution in crevices formed on SSs and Ni-based alloys tend to contain high concentrations of Cr3+, Fe2+, Ni2+, and CI" ions. The hydrolysis reactions of these cations normally take place inside the crevice lead to a decrease in the pH causing an increase in the dissolution rate'-16'. The determination of CCS in simulated solutions containing the metal cations, makes the polarization experiments more complicated by the variaty of electrochemical reactions taking place. Consequently, the CCS determination was made using a series of increasingly aggressive solutions maintaining a relation between an increasing chloride ion concentration and a decreasing pH. This was achieved by forming a solution of different molarities of NaCl where the pH is adjusted by HC1 acidification^10]. This concept, taking into account the chloride ion concentration the pH, was followed in the procedure of determination of the composition of the critical crevice solution. Results of potentiodynamic tests, in terms of critical pH (Figs. 2-10), showed that no direct relationships between temperature and chloride or pH as defined by Crolet'91. -226- Applying this criterion to the present findings indicate that the critical pH values lie between 1 and 1.5 (table 5). On the other hand, these findings seem to be generally in line with the Oldfield and Sutton'101 concept in the sence that increasing both chloride ion concentration and temperature lead to an increase in the peak current which in turn, indicative of the dpH. At 25°C, dpH for Inconel-625 alloy is zero, while the C-276 alloy remains resistant to film breakdown at this pH. This is in agreement with Russel^ who found that the dpH of alloy C-22 (developed from the original alloy C by varying the Mo content of the alloy) is less than zero at 20°C. The critical pH value for crevice corrosion can be considered an indicator to the susceptibility of an alloy to crevice corrosion, and thus can be utilized in the ranking of materials based on their resistance to this type of attack. The critical pH values determined by Oldfield and Sutton^11' for SSs 304L, 316L and 904L are 2.2, 1.65 and 1.25, respectively. The effect of increasing of both chloride ion concentration and temperature can be seen by comparing Figs. (2-10) and Tables (2,3 and 4). Both the peak-and passive currents increase with an increase in chloride ion conentration and temperature. According to the data, the critical pH values defined by both methods have been increased. This Phenomenon has been verified by Rosenfeld1171 who studied crevice corrosion of some SSs alloys. He found that the alloys were activated at higher pH values as temperature increase. For example, the lOCr lONi Ti steel was activated at pH values of 3 and 1.4 at 100°C and 20'C, respectively. This also appear for steel 18Cr 12Ni 3Mo Ti which was activated at 2.4 and 0.34. The results reflect the effect of temperature on dpH as well as the role of Mo in improving the resistance to passvity breakdown. -227- Si I .8: I I t O O Q O 8 8 8 gun / yu * •?•<* | -228- -229- m^^MMM^^Mmsm*^™^nr 10 90 90 »o CO TO I 0 0 i •foooot : i :r;~-~'.-:... ,•: :,:. •_' ' •~^'- - '••-'.: :..-•:•: : '.'i-.: : • - ']-?:: •• :•:'•'.' ooooz i •'•' "• ••• '•'" "-': :: 1 . •.,--•/ '.-'•'•'.'\.'-': '"' X o ^j-ooooe - CO i •K'.:::' •*-oooor • 1 I 1 00009 3 t i 00009 ; I j \oe.z-om\ I ' 9i9 • ID' i -oooot I 1 . .J 1 J ' y- 00008 7" M SI X PI: ''J 4-1 I'- * zuo / yu * tt*d | -231- zuo / yu ' tMd | -232 ju» / yu' tied | -233- &&1 '"•"p- f "T i i 1 pBffi yHEj •JL-. u-K' fvfa Kftl H| I?v« KW-{jCi |j X «ify a tatf NB* Fgii fl§ I [ffid (©3 IgL |j pg-'i t'C' TT -«j i'&'! f*S fijv: !v > A—A- A- oil ::f,;'-, H-:; -234- 1 > • • * it. ' I la S ! : i : i ; 1 CI i : ! ; ! B • i ' . : i I . J9. !5!i •. • •!:'• • f* *: • •:•. \\r.z\i : s.-;1 ;SliI.. I • . '"; . 5' •« a -.'•IS:'- ';; '.• '•'•;'•< Si !i;L;. \ ... .. "•'•-!. ..' • i" •••;•; 1 i:i::;" ;::. -r^r- •"M^j ll Si.- —^ =^ 1 B• • • • >! V-A \ \ V-A-A V- V 8 jura / V" ' *•• -235- Table (5): Critical crevice solution pH NaCl Critical pH Temperature Concentration Critical pH OldpeldA (C) (M) Crolet Criteria Sutton Criteria 25 1.5 0 0.5 50 1 0.7 100 1.0-1.5 0.8 25 1 0.5 1 50 1.5 1 100 0.7-1.5 1.8 25 1 0.7 4 50 1 1.2 100 1.5 1.7 -236- Passive current: During the induction period preceeding the breakdown of passivity, the driving force for causing changes in the crevice solution composition is the flow of metal ions across the metal interface. This is termed the passive current Ipass, although it may have two components namely, the healing of film flaws by oxide growth and actual metal dissolution into the crevice solution. The magnitude of Ipass is important in determining the rate of acidification. It is an esential input to the mathematical model to predict time to breakdown111' . The Ipass is determined by employing a solution similar to that inside a crevice, in terms of chloride ion concentration and pH. It is clear from Figs. (2-10) that the values of Ipass for alloy C-276 is generally lower than the corresponding values for Inconel-625 at practically all tested parameters. For example at pH=0.0, the ratio Ipass Inconel-625/Ipass C-276 ranges from 2 to 10 under similar chloride ion concentration and temperature conditions. The superiority of alloy C-276 is displayed when comparing the response of Ipass to pH. For example at 4M NaCl the ratio of Ipass (pH = 0.0)/ Ipass (pH=0.7) is about 2000 and 1000 at 25 and 100°C, while for inconel-625 alloy this ratio is 30 and 100, respectively. From the above results it is clear that alloy C-276 is more resistant to pH variation than Inconel-625 alloy. 5. Conclusions 1. The value of the critical pH is strongly dependent on both the electrochemical method and criterion used to evaluate the data. The two methods used in this study produced results that differ significantly. 2. The dpH for Inconel-625 alloy is 0.0 at 25'C and 0.5 M NaCl while not detected for alloy C-276. .'3. dpH increases with chloride ion concentration and temperature. -237- 4. Alloy C-276 is resistant to crevice corrosion thnn Inconcl-625 alloy under similar experimental conditions. 5. References 1. W.Z. Friend. Corrosion of Nickel and Nickel alloys, John Wiley and Sons New York (1980). 2. GA. Crangnolino and N. Sridhar. Corrosion. Vol. 47, No.6(1991).. 3. Corrosion and Corrosion Handbook. Edited by Philip A. Schewlitzer, P.E. consultant J Chester, New Jersey. Second Edition Copyright (1989). 4. R.M. Kain. Corrosion 186 paper NO. 229, National Association of Corrosion Engineers. Houston, Texas (1986). 5. J. Postlethwaite. Canadian Metallurgical Quarterly Vol. 39, (1983). 6. ZJ. Zuo, Z. Jinel, R. Sun, Y. Xu and X. Feng. Corrosion. Vol. 44, (1988). 7. D.W. Siitari andR.C. Alkire. J. Electorchem. Soc. Vol. 129 No. 3 (1982). 8. P.H. Cannington and R.J. Taylor. Corrosion Australasia Vol. 15 No. 3 (1991). 9. J.L. Crolet, L. Seraphlin and R. Tricot. Revue Metallurgei. Vol. 72 (1975). 10. J.W. Oldfield and H.W. Sutton. Br. Corros. J. Vol. 13 (1978). 11. J. W. Oldfield. International Materials Reviews. Vol. 32 (1987). 12. F.P. Ijsseling. Br. Corros. J. Vol. 15 (1980). 13. S.M. Sharland. Corros. Sci. Vol. 33 (1992). 14. M.K. Watson and J. Postlethwaite. Corros. Sci. Vol. 32 (1991). 15. W. John and S.K. Chan. Corrosion. Vol. 40 (1984). 16. J.W. Old field and W.H. Sutton. Br. Corros. J. Vol. 15 (1980). 17.1. L. Rosen feld. Localized Corrosion. Published by NACE Houston, Texas (1974). Editors R.W. Stachle, B. F. Brown and A. Agrawal. -238- EG rso^s//^ First International Spring School & Symposium in Advances in Material Science (SAMS 94) Cairo, 15-20 March 1994 CREVICE CORROSION OF HASTELLOY C-276 AND INCONEL-625 ALLOYS IN CHLORIDE ENVIRONMENT II: Effect of Bulk Solution Environment M.M. A- Gad, AJL El-Sayed and S.M. El-Raghy* Dept. ofMetalliurgy, Nuclear Research, Centre, Atomic Energy Authority, Cairo, Egypt *Cairo University, Faculty of Engineering Abstract Electrochemical techniques have been applied to study the crevice corrosion resistance of two nickel-based alloys *namely, Hastelloy C-276 and Inconel-625 in acidified sodium chloride aerated solution (pH=3), in the temperature range 25-100'C. Cyclic polarization (where the potential is reversed at 1.0 V(SCE) was performed on the two alloys at 25,50 and 100°C in the same pH chloride solution. Both alloys were resistant to this type of attack at 25°C, whereas only alloy 625 suffered crevice corrosion at 50°C, and both alloys were severely attacked at 100°C. This was clearly reflected in the value of AI in the hystersis loop and also in the values of both Ecorr and Eb. The elechtrochemical results were substantiated by SEM investigation. The findings were interpreted in terms of the effect of chromium and molydenum contents on the passive film characteristics and the influence of pH during polarization. Introduction Nickel-based alloys that have both high chromium and molydenum contents are resistant to corrosion in a wide range of oxidizing and reducing environments. They exhibit -239- good resistance to uniform and localized corrosion in chloride solutions and to stress corrosion cracking (SCC). The resistance to SCC is attibuted to high nickel contend . Hastelloy alloy C-276 and Inconel alloy 1-625 are normally used in the construction of radioactive waste storage containers which are employed in deep burial systems . These passive metals rely on the presence of a thin oxide film for their resistance to corrision. Halide ions, especially chloride ions, can partially or completely destroy such films leading to pitting corrosion on freely exposed surfaces and to crevice corrosion in areas of restricted access for dissolved oxygen^ . Inconel alloy 1-625 has superior mechanical properties because of the matrix stiffening effect provided by the addition of niobium and molybenum to the nickel-chromium matrix^ . Alloy C-276 has superior resistance to localized corrosion by the virtue of its higher Mo and W content. Alloy C-276 was developed from the original alloy C (higher carbon and silicon) by reducing the C and Si content to produce an alloy less susceptible to the precipitation of M6 C carbide and the intermetallic \i phase during welding. Alloys C-4 was further0^ developed by reducing Fe. Mn and W content which leads to further reduction in the u phase. At present, alloy C-4 is preferred as a construction Tiaterial for nuclear fuel Ffil waste disposal containers in Europe1 . Despite their excellent corrosion resistance in chloride solutions and brines. Ni-Cr-Mo alloys may suffer from uniform corrosion, localized corrosion (crevice and pitting) and stress corrosion cracking at elevated temperatures^1, K Other factors affecting their resistance to corrosion include: aeration and presence of other oxidizing agents (e.g. Fe , H202 and CIO") pH residual or applied stresses, previous thermal history, chloride concentration^ and the presence of tight crevices' '' . The hydrolysis of salts such as MgCL leading to lower pH and water radiolysis resulting in the formation of oxidizing agents such as 02, H20, and CIO" must -240- also be taken into consideration when these alloys are used in nuclear waste disposal applications. The present work aims at the investigation of crevice corrosion of alloys C-276 and 1-625 in chloride containing low pH solutions in the temperature range 25-100°C. Experimental Crevice corrosion studies were conduted on solution annealed Hastelloy alloy C-276 and Inconel alloy 1-625. Test specimens were in the form of discs (12mm diameter and 3mm thickness) with nominal composition given in table (1). The crevice corrosion bulk environment were acidified 4M NaCI (pH=3). An artificial crevice was created between the flat polished specimen (600 grit finish) and a (PTPE-Teflon)disc (12 mm diameter) in a special specimen holder (Fig. 1). The crevice assembly was located in a conventional multinecked 1 liter glass corrosion cell fitted with an external saturated Calomel electrode the electrolytic connection being a salt bridge to a luggen probe. An BG&G PARC Model 350A Electrochemical Corrosion System was used for cyclic potentiodynamic polarization scans. Specimen surface was examined after test using Scanning Electron Micro-scope (SEM) (Jeol JSM T200). Results Cyclic potentiodynamic polarization test starting at Ecorr (scan rate = 0.5 mV/sec.) at 25,50 and 100'C showed that both alloys exhibited a conventional hysteresis loop (Figs. 2, 3). The electrochemical parameters characterizing the passivation breakdown sometimes called damaging potential (E,) on the upward scan and repassivation termed the protection potential (E ) on the downward scan could be detected. Figure 2 and 3 showed that at 25°C both alloys exhibited resistance to crevice corrosion. However, for alloy -241- -O fllHG T«T SPECIMEN D WASHER SAMPLE HOLDER PLUG S*MPLE HOLOEH 60DT SAMPLE HOLDER CAP Figure 1 : Crevice Assembly Used -242- irf 102 103 10A 10s 106 107 nA/cm* to2 io3 10* id5 M>fr to7 to8 nA/cm uoo r-100'c uoc 1 i Figure 2 : Cyclic Potentiodynamic Polarization Curves for Inconel-625 in 24% NaCl Solution ,pH=3 -243- 1200 25'C 0800 /& ^ 0.400 1 0000 XJ1 -0.400 J^. 10° id 102 103 104 I05 I06 107 nA/cm2 1200 S0*C 0*00 f 20(00 A £ j 1 0000 ~~<.t -0400 3/ 10° W1 10* 103 10* 105 106 107 nA/tm? «*>!• IOO'C 0800 = 040C 0000 -OAOO d2 ^3 S 6 7 io° JO' IO IO JU^'IO IO IO nA/ccn* Figure 3 Cyclic Potentiodynanic Polarization Curves for Hastelloy c-276 in 24% NaCl Solution ,pH=3 244 625 the scan indicates that corrosion may take place. SEM (Fig. 4) reveals that this is not a crevice type corrosion. For alloy C-276 (Fig. 3) it is clear from the backward scan that the alloy remains active down to about 0.7V, SCE. Whereafter the alloy exhibited a passive trend. It is clear from the diagram that EC0JT and E for alloy 625 at 25°C are 0.00 and 0.07 V,SCE, respectively, while for alloy C-276 the value of E is 0.2 V.SCE. Both values and others at higher temperatures are summarized in table (2). At 50°C the behaviour of the two alloys is completely different (Fig. 2b and Fig. 3b). For alloy C-276 the scan is almost similar to that at 25 C except that there is a considerable difference (about 0.2 V) in the value of Enn„ SEM examination (Fig. 5-C) of the surface under the crevice indicates the beginning of some tendency to active dissolution. However, for alloy 625 the hysteresis loop indicates that a severe damage of the passive film has occurred. SEM examination of both free surface and creviced area confirms the cyclic potentiodynamic findings (Fig. a and b). Tests at 100°C showed that both alloys suffered crevice corrosion (Figs 2 and 3). The values of Ecorr and E for alloy C-276 are -0.25 and -0.2 V.SCE respectively. These values are different from those obtained at 25 and 50°C (table 2). The hysteresis loops for both alloys indicate that the passive film is completely broken down. SEM examination of both the creviced and uncreviced surfaces substantiates the polarization findings. SEM micrographs shows that active sites are initiated under the crevice former and may be extended to the crevice mouth (Figs 6 and 7). Discussion The crevice corrosion behaviour of the nickel-based alloys Inconel-625 and Hastelloy C-627 in low pH, high chloride environment in the temperature range 25-lOO°C showed a distinct difference. At 25 and 50°C alloy C-276 -245- [a] Inconel-625 Alloy [b] Hastelloy C-276 Alloy Fagure (4): SEM micrographs in 24% NaCl Solution at 25 c (1000X) - 246 - [a] Inconel-625 Alloy [c] Hastelloy C-276 Alloy Figure 5 SEM micrographs in 243 NaCl Solution at 50 C (lOOOx) -247- [a] Inconel-625 Alloy > Figure 6 : SEM micrographs in 24% NaCl Solution at 100 C (lOOOx) - 248 - [a] [c] [b] Hastelloy C-276 Alloy Figure 7 : SEM micrographs in 24% NaCl Solution at 100 C (lOOOx) .249- showed a remarkable resistance to crevice corrosion compared to alloy 1-625. However, at 100°C both alloys suffered severe crevice corrosion. Discussion of the present results shall be done in terms of the alloy performance in relation to both environmental, electrochemical parameters and the alloy composition. Effect ofpH on the corrosion behaviour ofNi-alloys: It is reported in the literature^ 1()1 that the corrosion resistance of nickel and its alloys depends on the pH and presence of the oxidizing agents in the corrosion medium. They have the tendency to form a good and a resistant passive films in neutral or alkaline solutions free of oxidizing agents. These passive films may be damaged in acid or very alkaline solutions containing oxidizing agents' . It is clear that the ASTM standard test is conducted at extremely low pH values (pH = 1.6) compared to those used in the present study (pH=3). However, other studies' ' -3 using the Q brine (26.8% Mg CI2 +1 mol/Z NaCI+10 mol./l HCl) for the determination of the critical crevice temperature (CCT) for the alloy C-4 (mainly of the same composition as C-276 except some minor alloying additions like C and W), which have lower values than alloy C-276, found that this temperature is close to that obtained by ASTM test. The environmental conditions (pH and chlorinty) used in Koster et al'12' investigation are almost identical to those employed in the present investigation. Other studies141 on the two alloys. 1-625 and C-276. at pH=2 in chloride-sulphate solution showed that the CCT for both alloys are below the critical pitting temperature (CPT). Moreover, alloy C-276 displays a higher CPT and CCT (table 3). These findings are in good agreement with the results of the present work. Electrchemical parameters Table (2) shows that the values of corrosion potential -250- Ecorr becomes more active with temperature for both alloys. For alloy 1-625 the detection of E was possible at both 50 and 100 C. The superiority of alloy C-276 over alloy 1-625 is clearly demonstrated both 25 and 50°C. It is worth mentioning that E at values are known' 3' to depend on the scanning rate used during the polarization experiments and the values are relatively close to reality at extremely slow scan rates. It is clear that the breakdown potential values (E^), considered as an indication to the susceptibility to crevice corrosion^141 for alloy 1-625 are shifted into the active direction as temperature is increased (Fig. 2-a, b and c). Table (3) indicates that the breakdown potential at 100°C is 0.1V,SCE. However for alloy C-276 the Eb was only detected at 100°C. Although the Eb values for both alloys at 100°C are nearly the same, the value of Al (taken by some investigators as a measure of the desruption of the passive films)^15' is about 700 - fold less for alloy C-276 (Table 3). There has been a controversy over the meaning of the value of Eb in relation to the prevailing corrosion type. Szklarska181 found that for a given SS (e.g. 316, 304 and 321) the critical potential for crevice initiation is always more negative than the critical potential for pitting initiation, but there is no simple correlation between both. Also, she concluded that the crevice corrosion of SSs in NaCI solution starts inside of crevices in the form of pits and can therefore be considered as a special kind of pitting corrosion. The findings of Szklarska is in agreement with S.M.Roshdy^8' who found that for some steels, crevice corrosion is initiated at lower potentials than pitting and crevice is the prevailent form of corrosion at all test temperatures. On the otherhand, Asphani' ' uses the term damaging potential to describe the potential at which passivity is broken down rather than using the terms pitting or crevice potentials since it is difficult to distinguish between these two forms of attack with cyclic polarization tests. Moreover, Sridhar'191 found that the determination of -251- breakdown and protection potentials or the difference between them were not sufficient indicators to differentiate between types of localized corrosion. However, the present results of E. and E were derived from an experimental (crevice assembly) sec up in which only crevice type corrosion is allowed to take place. The values of Eb and E potentials then are those corresponding to film breakdown resulting from crevice corrosion. Effect ofMolybenum This investigation showed that alloy C-276 is more resistant to crevice corrosion than alloy 1-625. The major difference in alloy composition is the Mo content (table 1). There has been evidence in the literature^ * that Mo has a beneficial effect on the localized corrosion resistance of SSs and Ni-based alloys. The mechanism by which Mo may prevent crevice corrosion has been explained in terms of improving the resistance of passive film to breakdown and enhancing the repassivation rates^ ' . Other investigators have have detected very little Mo in the passive film on SSs and suspected the Mo effect on the film integrity or film breakdown^ '. Mo may also improve the resistance against crevice corrosion through the formation of a molybdate ions which is absorbed on the metal surface thus promoting passivation and reducing dissolution rates1241. A similar interpretation on the effect of Mo on the corrosion resistance of Ni-based alloys was given by Mulford et al^251 who studied the behaviour of the two alloys namely alloy 600 and alloy 625. The first is Mo free while the other contains about 9% Mo. Surface Morphology: Scanning electron micrographs for corroded specimens have shown that crevice corrosion sites occurred on both alloys tested at 100°C. These active sites were also detected at 50°C. for alloy 1-625 only, while alloy C-276 remained -252- unattacked. For alloy 1-625 at 40CTC, several active sites occured and they formed active area larger than that produced at 50"C. The active sites were formed beneath the crevice former or near the crevice mouth. However, alloy C-276 showed crevice corrosion at 100°C, the severity of attack for both creviced and uncreviced area is less than alloy 1-625 (Figs. 5 and 6). Conclusions 1- Nickel-based alloy C-276 is more resistant to crevice corrosion than alloy 1-625 at 25,50 C, in 4M NaCI(pH=3). both alloys suffered crevice corrosion at 100'C, where alloy 1-625 is relatively less resistant. 2- pH of the bulk environment has a marked effect on the crevice corrosion susceptibility of both alloys. 3- The superiority of alloy C-276 over alloy 1-625 may be attributed to higher Mo content. References 1. W.Z. Friend, Corrosion Of Nickel and Nickel - Alloys. John Wiely end Sons, Inc., New York (1980). 2. G.A. Cragnolino and N. Sridhar. Corrosion. Vol 47, No. 6. (1991). 3. J. Postlethwaite R.J. Scoular and M. H. Dobbin. Corrosion. Vol. 44, No. 4 (1988). 4. Corrosion and Corrosion Protection Handbook. Edited by Philip A. Schewlitzer, P.E. Consultant I Chester, New Jersey. Second Edition Copyright (1989) by Marcel Dekker, Inc. New York and Basel. 5. R.M. Kain. Corrosion / 86, paper No. 229 National Assocaition of Corrosion Engineers. Houston, Texas, (1986). 6". R.D. McGright. Conference on Scientific Basis for Nuclear Waste Mangement XIV, Boston, Massachusetts. Public Materials Research Society, Pennsylvania USA (1991). -253- 7. R.C. Ncirman, R. Robcrgc and R. Bandy Corrosion. Vol. 39(1983). 8. J. Post let hwaite. Canadian Metallurgical Quarterly. Vol. 22 No. 2 ( 1983). 9. ZJ. Zuo, Z. Jingle, R. Sun, Y. Xu and X. Feng. Corrosion. Vol.S(WSS). 10. J. Postlethwaite. Localized Corrosion of Alloys C-276 and 625 in Aacrated Sodium Chloride Solution At 25 to 500 C Atomic Energy of Canada Limited (AECL) Report No. 9952, Pinawa, Manitoba ROE 1L0 (1991). 11. M. Pourbaix. Atlas of Electrochemical Equilibria in Aquouse Solutions. National Association of Corrosion Engineers Houston, Texas, USA. Second English Edition (1974). 12. R.E. Schmitt and R. Roster, Electrochemical Corrosion studies on Metallic Packaging Materials for High Level Radioactive Waste. Kernforschungszentrum Karlsruhe GmbH, Karlsruhe Report KFK4039 (1986). 13. M. Pourbaix and E.D. Vermilionk. Corrosion. Vol. 27 (1979). 14. K.D. Efird and G.G. Moller, Mater. Performance Vol. 18, No. 7 (1979). 15. B.E. Wild. Corrosion. Vol. No. 8 (1972). 16. Z. Szklarska - Smilalowska and J. Mankowsk. Corro. Sci. vol. 18(1978). 17. S.M. Roshdy. J. Electrochem. Soc. India, Vol. 37 No. 9 (1988). 18. A. I. Asphahani. Mater. Perform, vol. 19, No. 12 (1980). 19. G.A. Cragnoli.no and N. Sridhar. Corrosion Vol. 47 No. 6 (1991). 20. J.N. Wnnklyn. Corros. Sci. Vol. 21 No. 3 (1981). 21. J. W. OUtfield. Corrosion. Vol. 46, No. 7 (1990). 22. R. M. Davidson and J. D. Redmond. Mater. Perforin. Vol. 27, No. 12 (1988). 2-1. R.L. Tapping. Corros. Sci. Vol. 25, No. 6(1985). 24. R.C. Newman. Corros. Sci. Vol. 25 (1985). -254- First International Spring School & Symposium in Advances in Material Science (SAMS 94) Cairo, 15-20 March 1994 A STUDY ON THE CORROSION BEHAVIOUR OF X-750 NICKEL BASE ALLOY MJL EI-Hady and ALA. Kassem Faculty of Petroleum & Mining Engineering, Suez Canal University, Suez, Egypt. Abstract The effect of chloride ion concentration on the anodic dissolution behaviour of X-750 alloy has been investigated using the potentiostatic anodic polarization technique. Increasing chloride ion concentration changes the type of attack from pitting to general corrosion. It was found that, the electrolyte temperature has a significant effect on corrosion behaviour. The effect of thermal aging treatment on corrosion behaviour was also studied. It was found that; the solution annealing at 1050°C for 2 hours followed by aging at 870*0 for 4 hours gave the best corrosion resistance in 1 N H2S04. 1. Introduction Alloy X-750 is a Ni-base alloy. It is used in varity of application in light water reactor (LWRS) as core bolts, Fuel assembly, hold down springs, guide pins, control rod drive seals and jet pump beams. It suffers from the stress corrosion cracking (SCO problems which affect the safety of the nuclear power station and increases costs for repair. The factors affecting the SCC of alloy x-750 have been reviewed'1'. These factors were defined; heat treatment (single, or double aging), chromium and zirconium additives, the applied potential, the strain rates and the contents of sulphate and oxygen in water. Intergranular stress corrosion cracking (IGSCC) has been studied in two similar - base -255- alloys, alloys 600 and X-750 in deaerated steam at elevated temperature (380°C)121. The rate of IGSCC is much faster in both alloys when they are heat treated so that carbon is kept in solid solution instead of precipitating as grain boundary carbides. The microstructural changes and grain boundary chemistry of high strength age hardenable inconel X-750 have been studied using electron and Auger microscopy following a sequence of thermal treatments in the carbide precipitation temperature zone of 704°C to 871°C^3l Austentic X-750 alloy constitutes mainly the y matrix, y (Ni3 Al and Ni3 Ti) intermetallic prcipitates inside the grains together with MC type (TiC, NbC) and M23 C6 type (Cr23 C6) carbide. The effect of chloride ion concentration on the anodic dissolution of nickel has been investigated using the potentiostatic anodic polarization technique and metallographic observations^. The aim of the present work was to study the effect of thermal aging treatment, chloride ion concentration and electrolyte temperature on the corrosion resistance of X-750 alloy. 2. Material and Experimental Procedures Samples of the alloy X-750 for various experiments in this study were prepared. The chemical composition of the used alloy is given in the following table, (wt %) Ni Cr Fe Ti Al Nb+Ta C >75 14-17 5-9 2.25-2.75 0.4-1 0.7-1.2 < 0.08 Mn Si S Co Cu < 1.0 < 0.5 < 0.01 < 0.2 < 0.3 The samples were solution annealed at 1050°C for 2 hours followed by water quenching. A part of he samples were aged at 870"C for 4 hours and another part at 750°C for 48 hours. The third part was exposed to douple aging by annealing at 816C/20 hours, AC + 704T/24 hours, (AC: air cooling). -256- Fig- <1> shows the microstructuros of different heat treated samples which were used in this work (solution annealed at 1050 C for 2 hours, solution annealed at 1050 C for 2 hours, then annealed at 750"C for 48 hours, solution annealed at 1050°C for 2 hours, then annealed at 870"C for 4 hours, and sample solution annealed at 1050CC for 2 hours followed by douple aging, annealed at 816°C/20 hours, AC + 704X/24 hours, AC). The corrosion behaviour of these samples was studied in IN H9 S04 and IN H2 S04 with different percentages of IN NaCl (2,4,10,20,30 and 50%). The potential-current measurements in the anodic direction were carried out with potentiostatic regulator (wenking potentioscan Pos 73) using quasi stationary technique. 3. Results & Discussion 1. Effect of heat treatment on the corrosion resistance of X-750 Ni-base alloy in IN H2S04 In order to evaluate the effect of heat treatment on the corrosion behaviour of this alloy, anodic polarization curves for sevral specimens with different heat treatment conditions were studied. The obtained results are compared in Fig. (2). Table (1) shows the effect of heat treatment on the corrosion characteristics of X-750 Ni base alloy in IN H2S04. Table (1) : Effect of Heat Treatment on The corrosion Characterestics of X-750 alloy in 1 H H9S04 Passive Specimen Condition ranee (mV) (mA) 1. As received 200-1000 350-550 2. Solution annealed at 1050 C Cor 2 horns 400-1000 250 ."!. Sample solution annealed at 1050 C tor 2 h«uirs. 200-1000 100-150 then annealed at 750 (.' for IS hours. •I. Samples solution annealed at 1*>50 C tor 2 hours 20O-!»IM) :'il)-t;i» followed h\ annealinj; at S7() (' lor 4 hours. 5. Samples solution anneald at 1050 C lor 2 hours I (o-Siiu l'i-22 tallowed l)v douple aying (annealing at MB (.720 hours. AC + 704 C/2-J hours. A( •>. -257- From Fig. (2) and Table (1) it could be stated that, the specimens which has the best corrosion resistance is the one which had solution annealing at 816°C for 20 hours, AC + 704°C for 24 hours, AC, (double aging), passive range of 140-800 mv and passivation current of 18MA. Also the samples solution annealed at 1050°C for 2 hours followed by annealing at 870°C for 4 hours has very good corrosion resistance (passive range of 200-900 mV and average passivation current of 55 mA). The effect of heat treatment on the corrosion characteristics of X-750 Ni-base alloy can by explained as follows. A heat treatment that produces or can not completely dissolve C-rich M23 Cg type carbide precipitation along the grain boundary results in the highest corrosion rate. (Fig. 2). Aging at 870*0 for 4 hours, initiated the dissolution of the grain boundary Cr-rich, M23 Cg type carbides with complete dissolution of these carbides (Fig. 1-b). The heat treatment condition of the two stage treatment (Double aging) results in the formation of incoherent y particles and secondary broken discrete carbides at the grain boundaries (Fig. 1-c). 2. Effect of Chloride ion concentration on the anodic dissolution behaviour of X-750 Ni-base alloy The inhibition of corrosion in most metals is due to the presence of a passive or protective film. The formation of the passive film proceeds by the anodic dissolution of the metal. The presence of certain anions, particulary chloride, destroys passivity and leads to localized corrosion, which takes the form of pitting. Fig. (3) and Table (2) shows the effect of choride concentration on the corrosion characteristics of X-750 Ni-base alloy which are solution annealed at 1050°C for 2 hours followed by water quenching, then aged at 871°C for 4 hours followed by air cooling. -258- Fig. (1) : The microstructure ofX-750 Ni-base alloy. a) Solution annealed at 1050VI2hrs + 750 I48hrs. (X150) b) Solution annealed at 1050 V/2hr + 870 CI4hrs. (X250) c) Solution annealed at 1050V12hrs + double ageing (X250) -259- a /Vii.vl lO(J0 . 800 rnvGw) ton 200 CtOU Lou I t,nm /CM •) FI'A'. T2; ; Anodic polarization curves of X-750 nickel base alloy with different heat treatment conditions in IN H.JSOJ -260- •JOII L_J I L O.I I 10 Unj I IttlA/l!)) > Fig' (3): Anodi polarization curves ofX-750 nickel base alloy aged at 870 V for 4 hours in different solution mixtures oflNNaCl + 1NH2S04. -261- The testing electrolyt composition was 2,4,10,20,30, and 5% of IN NaCl in IN H2S04 at room temperature. Table (2) : Effect of Heat Treatment on The Corrosion Characterestics of X-750 in HS04-NaCl Mixture. Specimen Passive Condition Environment range IP (mV) (mA) Sample solution 2%NaCl+98%H2S04 120-900 Rangec annealed at 4%NaCl+96%H2S04 100-900 from30 1050°C for 2 10%NaCl+90%H2SO4 100-800 AM to hours then 20%NaCl+80%H2SO4 100-900 50 MA annealed at 30%NaCl+70%H2SO4 0 -850 100 870'C for 4 hours 50%NaCl+50%H2SO4 0 -600 120 Table (2) shows the effect of chloride ion concentration on the corrosion characteristics of X-750 Ni-base alloy. From Fig. (3) and Table (2) it was found that, samples which are aged at 870°C for 4 hours can withstand the chloride ions concentration up to 20%. When the chloride ions concentration increases above this percentage, the corrosion resistance significantly decreases. This means that, a critical concentration of the aggressive ions exists (20%), up which the break down of the passive film is localized resulting on pitting. Beyond this concentration, the area of attack increases given rise to general corrosion. 3. Effect of temperature on the corrosion characteristics of X-750 alloy Fig. (4) shows the effect of temperature on the corrosion characteristics of X-750 for samples solution annealed at 1050°C for 2 hours then aged at 870°C for 4 hours; the testing electrolyte was 10% IN NaCl + 90% IN H2S04. -262- f inno 1 ( * 20 C t * 50 C 1 acio 1 • flu c; «oo 4 III I 200 0 00 2(10 0. I IO ICO lull i (iiiA/irm, Fig. (4) : Anodic polarization curves ofX-750 nickel, base alloy aged at 870 V for 4 hours in solution of 10'/, In NaCl + 90% IN H2S04 at different temperatures. -263- From the figure we found that the passivation range for three testing temperatures (20,40,80°C) is from 100 to 900 mV and I values are 50,100,1100 mA respectively, which means that, temperature has a great effect on the corrosion behaviour of X-750 Ni-base alloy in 10% IN NaCl + 90% IN H2S04. 4. Conclusions 1. Samples which are solution annealed at 1050"C for 2 hours followed by water quenching, then aged at 870°C for 4 hours and the samples which are exposed to double aging by annealing at 816*C/ 20 hours, AC + 704'C/24 hours, AC, has the best corrosion characteristics. 2. A critical concentration of the aggressive ion exists (20%). at which the break down of the passive film takes place on X-750 Ni-base alloy resulting in pitting. Beyond this concentration, the area of attack increases giving rise to general corrosion. 3. The temperature of corrosion attack has a very great effect in the presence of chloride ions. 5. References /. S. W. Sharkawy and F.H. Hammad, Failure problems of Alloy X-750 in Light water Reactors. 5th Conf. of Nucl. Sc. & Appl.; Vol.1 162-173, 1992. 2. Yulin shen and Paul G. Shewmon. Intergranular stress corrosion cracking of alloy 600 and X-750 in high-temperature deaerated water stem. Metallurgical. Transactions A. Vol. 22A August 199 l-p. 1857. •i. B. Mishra, A.K. Sinha and J.J. Moore. Effect of Single. Aging on MIcrostructure and impact property of inconcl x-750, Metallurgical transaction A, Volume 16 A. May 1985 P. 821. 4. M. Zainin and M.D. Ives. Effect of chloride ion concentration on the Anodic Dissolution Behaviour of Nickel, corrosion -Nace, Vol. 29, Mo. 8, August, 1973. -264- ECfSo 4SAL First International Spring School & Symposium on Advances in Materials Science (SAMS 941 15-20 March 1994 SULFIDATION OF IRON-25% MANGANESE ALLOY IN H2S-H2 ENVIRONMENT AT TEMPERATURE IN THE RANGE 973-1173K FA Elrefaie, HA. Ahmed, N.A. Abdd Rehelm' and W.W. Smeltzer* ABSTRACT The sulfidation behaviour of an Fe-25%Mn (by weight) alloy has been studied in H2S-H2 J, environment at 1C 1 Department of Metallurgical Engins; ling. Faculty of Enoini-.-ing, Cairo University, Giza. Egypt. 2 Department of Materials Science and Engineering, and Institute of Matenals Research. McMaster University, Hamilton, Ontario, Canada LBS -265- INTRODUCTION Investigation on the sulfidation behavior of manganese [1], molybdenum [2], and vanadium [3] have shown that these metals exhibit higher sulfidation resistance as compared to the resistance to sulfur-bearing atmospheres exhibited by iron, nickel, cobalt, chromium and alloys based on these met.i'- [4,5]. The sulfidation behavior of an iron alloy close to its molybdenum saturation in H2S-H2 atmospheres, 10~4 EXPERIMENTAL TECHNIQUES The Fe-25% Mn alloy was prepared by melting iron and manganese chips of 99.97wt% purity. The iron purity is about 99.99% (Marz grade, Materials Research Cooperation). Cylindrical rods, 7.6 cm long and 1.9 cm diameter, were initially produced by electric-arc melting of iron and manganese under argen gettered of oxygen by titanium. Samples as disks 0.1 cm thick were subjected to metallographic polishing finishing with 1 micron diamond paste. These specimens were then sulfidized in flowing H2S-H2 atmospheres 10"4 Marker studies to understzi -1 the sulfidation mechanism is carried out by placing Pt wire, 25 micron in diameter, around the specimen which then sulfidized at Ps2 = 10 Pa and 1073K. RESULTS AND DISCUSSION The measurements of the sulfidation kinetics at 973K, 1073K and 1173K and different sulfur pressures are illustrated in Figs. (1) and (2). Figure (1) is a plot of the weight change per unit area, dw/a. versus time while Fig. (2) is a parabolic type of plot in which (dw/a)2 is plotted versus time Figure (2) indicates that the reaction kinetics is characterized by two -266- (a) ' K» Pi /^ ^^* ' -•- 1 Pi yS ^i^"^ " -*• lf-1 Pt yf J*^ -o w-a PI JS s^ ^—** :ao-' -*• «-j »> / jr^ -**^' -»- ie-« PI > 5 10 5 ^ so 100 150 200 350 300 T1#n« (mln) (0 PrailMft fl Svlliff / — 10 Pi -+- 1 Pi -*• ll-l Pt 20 • / / •*• «•» Pt £ u -*- v j «• O 15 -ft- 11-4 Pi < 10 I^-T -—-*- , : »—: 150 200 Tim* (mln) J Figure (I); Sulfidation kinetics of Fe-25% Mn alloy in H,S-H2 enviroment at 10 < Ps,(Pa)<10: a) at 973 K, b) at 1073 K, and c) at 1173 K -267- no we 300 au TWK* (nikO •00 400 300 0«** 100 300 400 MO TtKM Ifflln) (c) ?»».•«»• »t Bwltw — 10 Pa (00 -+- 1 r. •*• lt-1 P. -»• if•» r« -*• *-J f• I -*- *-« *• 8' too 5 • / / / -/° • Iff. so 140 300 ISO 300 350 Figure (2); Parabolic kinetics of Fe-25% Mn alloy in H.S-Hj enviroment at 10U < PSj(Pa)<10: a) at 973 K, b) at 1073 K, and c) at 1173 K -268- stages, the first is an initial transient stage which ranges from 13min at 1173K and Ps2=lOPa to 95min at 973K and PS2=iOPa. The second stage is parabolic, the parabolic rate constants obtained from the slopes of the plots for this stage is given in Table (1). Figure (3) is an optical micrograph that shows the cross section of scale formed on the Fe- Mn alloy at 973K and 1tr* From Figs. (4) and (5), it can he seen that the solid reaction products formed by sulfication at 1073K and 1173K consists of three sublayers similar to those formed at 973K overlaying internal precipitation zone, the thickness of the internal precipitation zone increases as the 4 3 pressure decreases, except at PS2=1CT Pa at 1073K and P$2<10" Pa at 1173K where only one layer is formed overlaying the internal precipitation zone. Figure (6) is a scanning electron micrograph for fractured cross section of sulfide layer formed at 1073K and PS2=10Pa, this figure demonstrates-the formation of an outer columnar layer. Figures (7), (8) and (9) demonstrate the reflection X-ray diffraction patterns obtained for scales formed at 973K. 1073K and 1173K. The inner two sublayers are composed of alpha-MnS for all three temperatures, while the outer layer is composed of pyrhotite (FeS). 4 < n 3 At PS2=10" Pa at 1073K. and PS2 1 " Pa at 1173K only one layer is formed and is composed of alpha-MnS. The diffraction patterns of the single layered-scale samples show diffraction lines of alpha-Fe which indicate the transformation of gamma-Fe to alpha-Fe at scale/alloy interface. Electron-probe scans, Figs. (10). (11) and (12) give the Fe and Mn profile along the cross section of the sulfidized samples. It indicates the following: 3 (1) At 973K and Ps2=10" - 10Pa (Fig. 10), the scale is layered in the sequence manganese sulfide (inner layer)/iron sulfide (outer layer). 1 (2) At 1073K and PS2= 0Pa the structural sequence is manganese sulfide (inner layer)/iron sulfide (outer layer). This changes as ir.n pressure gets smaller; at PS2=10"^ and 10"2 Pa the layers order is as following: internal precipitation zone/inner layer composed of manganese sulfide/outer layer composed of iron 3 sulfide The thickness of the internal precipitation zone is larger at PS2=10" Pa. At -269- P , = I Pa Ps2 = 10 Pa s2 P = I0*3 Pa sZ Ps2 = ,0 Pa Fig. (3) : Microstructure .'••-itures of the three layered scale formed at 973K and different P ,. -270- Ps2 = 10 Pa P., = I Pa *&j*fc 3 P = 10"* Pa Ps2 = 10" Pa S2 Fig. «f) : Microstructure features of the scale formed at l073Kand different Ps2- - 271- Ps2=in-4 Pa only a single layer composed of manganese sulfide is formed over the internal precipitation zone. 3 4 (3) At 1173K a single layer is formed at PS2=10" and 10" Pa overlaying the internal precipitation zone, while at higher pressure the layer sequence is as following: internal precipitation zone/manganese suifide/iron sulfide. The sulfide precipitates were manganese sulfide containing a smaller iron concentration than in the scale phase. These results of the electron-probe scans agrees with X-ray diffraction results and the metallographic examination. The parabolic sulfidation rate constants of the alloy are compared to the parabolic sulfidation rate constants of pu>v manganese and pure iron in Table (1). It can be concluded that in the case of formation of three-layered sulfide at the three investigated temperatures, the growth of the sulfide layer on the Fe-25% Mn alloy is of slower rate than the growth of iron sulfide on pure iron and of a faster rate than the growth _4 of manganese sulfide on pure manganese. A single layer is formed at PS2=10 Pa and 1073K, and at Ps2<10'3Pa and 1173K; at these conditions the sulfidation rate constant of both the alloy and pure manganese are almost the same, while the sulfidation rate constant of pure iron is zero. It is clear that the parabolic rate -.onstant increases as the sulfur pressure increases. Arrhenius plots of the parabolic rate constant at several sulfur pressures are given in Fig. (13); the activation energies calculated from the slopes are recorded in Table (2). An average value of 11,000 ±1500 cal/g. atom is estimated under a constant sulfur pressure within the investigated sulfur pressure range of 10"2 to 10Pa. This Is to be compared with a value of 12,000 cal/g. atom for the sulfidation of iron and a value of 21,000 ±2000 cal/g. mole for the growth of manganese si'lfide on Mn under constant suifur pressure [9]. Figure (14) gives the results of the marker investigation; :r>° figure indicates very clearly that after sulfidation the marker is located at the interface of the columnar MnS/equiaxed MnS sublayer. All these findings are consistent with the consideration INat the outer FeS and MnS columnar sublayers grow by outward migration of cations v/hile the inner equiaxed MnS sublayer and the internal precaution zone grow by inward migration of sulfur. -272- Table (1) : The parabolic sulfidation rate constants for growth of sulfide scales on Fe-25%Mn alloy obtained in this investigation and values calculated for the parabolic sulfidation rate constant for pure manganese (1) and pure iron (3). Units ofKp: (g2cnr4S-1) 973K 1073K 1173K s2(Pa) Fe-25% Fe- Fe-25% Fe Fe Mn Fe Mn Mn 2590 Mn Mn 10 9.8x10"8 7.3X10"7 2.2X10"7 1.3X10"6 3.5x10"9 3.9x10'7 1.9X10"7 9x10-9 5.2x10-9 1 6.4x10"8 5.4x10-7 1.3x10"7 8.5x10-7 2.2X10-9 1.8X10-7 1.1X10'7 5.1x10-9 10-1 3.6x10-fl 3.8X10'7 5.8x10-8 5.2x10-7 1.9x10-9 1.1x10"7 5.3x10'7 4.3x10-9 2.8x10-9 10-2 1.7x10"8 2.5x10"7 3x10-8 3.1x10"7 1.2x10-9 5.8x10-a 1.7X10"7 2.9x10-9 3 10- 1.1x10-8 1.4x10"7 9.7x10"9 1.2x10"7 7.6x10-'° 1.3x10-9 0 1.9x10-9 9 8 9 10-4 5.6x10- 6.4x10- 1.2x10- 0 1.3x10-10 0 ' Table 12): Activation energies for sulfication of Fe-25% Mn alloy at various H2S-^2- PS2(P*) Q.cal/g atom 10 9,410 1 10.729 10-1 10.610 10-2 12.610 -273- CONCLUSION The sulfidation behaviour of Fe-25%Mn allloy has been studied in H2S-H2 environment at 4 10- (1) The reaction kinetics are parabolic following an initial-transient stage that extends up to 95 minutes with increasing parabolic sulfidation rate. 4 3 (2) At 1073K and PS2=10- Pa, and at 1173K and PS2<10- Pa an external layer of manganese sulfide layer overlaying an internal precipitation zone is formed. At all other investigated conditions, an outer iron sulfide layer overlaying an inner manganese sulfide layer is formed. (3) The inner manganese layer consists of two sublayers; an outer compact columnar- sublayer and an inner equiaxed porous-sublayer. (4) Marker measurements indicate that the groMh of the inner equiaxed manganese sulfide sublayer is supported by inner migration of sulfur while the growth of the outer manganese sulfide sublayer and the iron sulfide layer results from outward mirgration of iron. -274- REFERENCES (1) F.A. Elrefaie and W.W. Smeltzer. Oxidation of Metals, vol. 16, Nos. 3/4, p. 267 (1981). (2) B.S. Lee and R.A. Rapp, J. Electrochem. Soc. vol. 131. p, 2998 (1984). (3) F.A. Elrefaie and W.W. Smeltzer. High Temperature Oxidation and Sulfidation Processes, ed. J.D. Embwy, CIM, Hamilton, Ontario, Canada, p. 323 (1990). (4) D.J. Young, Rev. High Temp. Materials, vol. 9, p. 299 (1980). (5) S. Mrowec and K. Przybylski, High Temp. Materials and Processes, vol. 6, p. 1 (1987). (6) F.A. Elrefaie and W.W. Smeltzer. Werkst. Korros., vol. 38, No. 9, p. 493. Sept. (1987). (7) F.A. Elrefaie, H.A. Ahmed, M.F. El-Demerdash and W.W. Smeltzer, J. of Engineering and Applied Science, Faculty of Engineering, Cairo University, vol. 39, No. 5. p. 1081, Oct 1992. (8) G. Romeo, W.W. Smeltzer and J.S. Kirkaldy. J. Electrochem. Sco„ vol, 118, p. 740 (1971). (9) E.M. Fryt, V.S. Bhide, W.W. Smeltzer and J.S. Kirkaldy, J. Electrochem. Soc, vol. 126, p. 673, 683 (1979). (10) G.G. Libowitz, "Proceedings of VII International Sympsoium on Reactivity of Solids", p. 17, Butterworth Publ., London '1972). Ps2 = 10 Pa ft.**• -CiCfcS£%.| :^T5 1 2 3 Pj2 = JO" Pa Pj2 - I0" Pa P$2 = 10" Pa .»l ^/f\ -. \- ... '• • N?S^> '•• Fig. (5) : Microstructure features of the scale formed at 1173K and different Ps2- Fig. (6) : Scanning electron micrograph of fractured cross-section of sulfide layer formed at I073K and P - = 10 Pa. -276- I P 2 (10 Pa) Pl2 (I Pa) Pi2(«o Pa) The upper surface ot The upper surface of The upper surface ol e 2nd layer the 2nd layer the 2nd layer c i a I S Sg S 1* l I I J_ s S I Ps2 (I0*' Pa) Pl2 (10"* Pa) Pl2 (10** Pa) The upper surface of The upper surface of The upper surface ol ihe 2nd layer the outer suliide layer (Fe) the 2nd layer c 5 •si =<0 S I —U n _1_ I ffiJLt l J. Fig. (7) i Reflection X-ray diffraction patterns obtained at T • *73K. - Ps, (10 Pa) c 5 Ps2 (I Pa) The upper surface of Pa (10"' Pa) The upper surface of the 2nd layer The upper surface of the 2nd layer c the 2nd layer c S c 5 c 2 1 L_, rk to -J oo 3 Ps2 do" Pa) Ps2 Ot Pa) 2 Upper surface of P5- (10" Pa) Upper layer Surface of the only 2nd layer layer formed (one layer) I • §~ & c c c Er.(a. a. a. Tl &ST 22 2 Fig. (S) : Reflection X-ray diffraction patterns obtained at ! . MnS s •q P(10Pa) Theuppersurfacofth j2 1 2ndlayer 1 *.\ sin Fig. (9):ReflectionX-raydiffractio n patternsobtainedatT=II73K. • * m 3 £ P(IO"Pa) -= Theuppersurfacofth s2 sulfide layer(onlyon lO 1 i MnS ayer) MnS 3 « 2I| f! 1 P,0Pa) The uppersuracofth 2nd layer 2 i— «g f \^ Distance Mn /"A",\ Distance Fig. (10) : Electron-microprobe scans along the cross sections of samples sulfidized at 973K and -3 Ps2 - 10 Pa; Ps2 I Pa (b); and P _ = 10 s2 Pa (c) -280- F±J\ Fe Mn\ r " UM'a Distance Distance Distance to (f) Fe (e) Mn Distance Distance Distance Pig. (ID : Electron microprobe scans along cross sections of samples sulfidized) at 1073 and Pj2 = 10 Pa (a); P^ = I Pa (b); 2 3 \\2 -- 10-' Pa (c); P$2 = | Distance Distance Fig. (12) : Electron-microprobe scans along cross sections of samples sulfidized at II73K and P$2 = 10 Pa (a)i P ? = I Pa (b)i Pj2 = 10*' Pa (c); 2 3 Pj2 = I0" Pa (d); and I0" Pa (e). -282- • P» • 10 P» •61 • — p» " t p# • -It -. • P»« 10«-1 P« . -6 5 o fit I0«-JP» I I" .7.4 .7.6 • •7.a •"— O.OOOM ooocar 0.00019 OOOMI 000093 000095 0.00097 0.00099 oootoi oooioi vtm Fig. (13) : Arrhenius plots ol parabolic rate contant at different Ps2- Fig. (I*' •' Optical micrograph shouing the position .if ft marker iti the sulfide senile. -283- First International Spring School & Symposium on Advances in Materials Science (SAMS 94) 15-20 March 1994 Effect of Ultrasonic Agitation on Chromium Electroplating S.M. Morsy*. A.Y. Hosny", M.W.El-Rafey and M.M.El-Sayed" * Metallurgy Dept. Nuclear Research Center. Atomic Energy Authority, Cairo, Egypt. " Department of Material Science, Institute of Graduate Studies and Research, Alexandria University. Abstract This research work presents the effect of ultrasonic agitation on chromium electroplating. Chromium electroplating divided into two methods,conventional plating bath with ultrasonic and without ultrasonic, and the second by using hull cell with ultrasonic and without ultrasonic. The results of experiments showed that ultrasonic agitation (30 Kilocycle/sec) increased the current efficiency deposition of chromium from the electroplating bath. The use of ultrasonic agitation in chromium electroplating permits production of good quality films on steel. The results of this work was confirmed by cyclic potentiodynamic polarization technique in 3% NaCl. I. Introduction In spite of a large number of investigations concerning the mechanism of the effect of ultrasonic field on the electrodeposition of metals, there is still no general theory of this process. Most investigators assume that the main effect of ultrasound is similar to mechanical mixing (1-4). At the same time, other investigators have ascertained that ultrasonic vibration induce a uniform composition of the electrolyte which cannot be attained by any other method used in electrodeposition of metal (5-10). Although the current efficiency of chromium plating is low, generally in the range of 10 to 25% for bright plate, and often nearer the lower than the upper range (11,12), a fairly high rate of deposition is achieved owing to the high current densities are higher than in most plating operations. Thus the generator or rectifier capacity required for chromium plating is higher than for most metal plating. With the ever increasing interest in high speed plating and the raising cost of metals, either an increase in the plating rate or a reduction in the metal ion concentration without affecting speed or deposit properties is industrially significant. The object of this research is to study, if ultrasonic agitation could help to achieve of these goals in chromium electroplating. -285- 11. Experimental procedure 11. I. Material Coils of Uddeholm Swedish stainless steel razor blades (22.4 mm width. 0.001 mm thickness! imported by Lord precision industries were used as substrate material. The chemical composition is given in Table (1) Table (1) Chemical Composition (% by weight) c Mn S Si P Cr 0.65 0.65 0.015 0.4 0.025 12.5 Steel strips were cut to desired lengths 18 cm and cleaned with trichloroethylene using ultrasonic cleaner for 2 minutes. Specimens were divided to give an area equal to 2.2 x 2.2 cm2 double face. The rest were wrapped by teflon tape, after chromium electroplating, the teflon tape was removed and specimens washed by distilled water and dried by warm air II.2. Chromium Electroplating Technique 112A. Conventional Plating Bath A 60 ml glass container was used as chromium electroplating cell. A fixed plexiglass holder was placed on the top of the cell to hold the steel cathodes and lead anodes of equal size vertically and to maintain a constant distance between electrodes. The cell was placed in an ultrasonic cleaner (30 Kc/sec) and the electrodes of the ceil were connected to DC. power supply unit. Chromium electroplating was carried out using the following conditions: Solution concentration 250 g/L Cr03 2.5 g/L H2S04 Temperature 45'C± 1C Plating time 5 minutes Current density 300 Ma/cm2 Voltage 5 volts -286- 11.2.2. Hull Cell The hull cell is the most generally used device for plating range tests. The celluse'1 had a solution of volume 267 ml. The anode of lead (7 cm x 55 cm x 0.15 cm) occupies the square cross sectional end of the cell, and the polished copper cathode (10 cm s 7.5 cm x 0.05 cm I is placed along the inclined side 11.3. Corrosion Measurement Electrochemical corrosion measurements were carried out using minipotentiostat, winking model (MP 87). Chromium coaled steel specimens were cleaned by distilled water using an ultrasonic cleaner and immersed in 3% NaCl. The potentiodynamic cyclic polarization was carried by a controlled potential scan beginning at -0.250 V (ECorr related) and extending to some vertex point, followed by a reverse scan back to the final potential. The following parameters were used throughout this work. Initial potential -0.250 V Final potential -0,200 V Vertex potential 0300 V Scan rate 1 mv / sec. Specimen 1.000 cm^ III. Results and Discussion The ratio between the weight of the metal deposited on the cathode to the theoretical weight as expressed in percent according to Faraday's Law is called the current efficiency. In chromium deposition one ampere / hour deposits 0.323 gm of chromium. The current efficiency is ideally 1007. for metal deposition, but in practice a value greater than 90% is regarded as satisfactory. l£ significant amounts of hydrogen are evolved, the efficiency decreased markedly and in the case of chromium could be as low as 10% (11. 12). The value of the cathode current efficiency is therefore an important economic consideration in an electroplating process. The main goal of this research to see if ultrasonic agitation could help to increase the value of current efficiency and quality of electrodeposition in the case of chromium electroplating. Figure (1) 6c (2) show that by using ultrasonic agitation of 30 kilocycle/second, and increase in both the weight of deposit and the current efficiency is observed after 10 minutes agitation. Figure (3) shows the relation between hardness and deposition time with the ultrasonic agitation and without it; a noticeable increase in hardness is observed in case of ultrasonic agitation, These results lead us to assume that the structure of chromium deposit undergoes greater -287- W<|»nmli;' :a / z 18 /. s- 12 6 J6'/*> < 6 8 to Tine (Mia.) Fig. (1) Relation betveen veight of deposit and depositioa time (t) vith ultra sonic (2) vithout ultrasonic Current efficiency * Time (Mia.) Fig. (2) Relation betveen current efficiency and depositioa tine (1) vith ultrasonic (2) vithout ultrasonic -288- Hardness Hv Fig. (3) Relation between Hardness and deposition time (1) ultrasonic agitation. (2) wMout ultrasonic agiution -289- change than that of chromium electroplating without ultrasonic agitation The propensity of a plating bath to give continuous or non- continuous coating is called covering power. This can be studied by using a Hull cell, Figure (4) shows that by using ultrasonic agitation, coatings of high covering power are produced, but only if the effect of agitation by ultrasound takes place at the beginning of ultrasonic agitation. As shown in table (2) the ultrasonic agitation has an effect only before the precipitation of chromium: after the deposition of chromium begins the ultrasonic field has no significant effect on the value of current efficiency Ultrasonic agitation can effect the nature of crystallization and the quality oi the chromium deposit. It is noted that the chromium deposits produced with ultrasonic agitation are less porous than deposits without it. This may be a result of the finer grain size of the deposits produced with ultrasonic agitation, as illustrated in figure (5) Corrosion resistance of the chromium deposits obtained with ultrasonic agitation and without was assessed by using the cyclic potentiodynamic technique. For each deposit, the test specimen potential was scanned from -0.250 V "Ecorr" ( Corrosion potential) related to an arbitrary anodic potential "Ey" and back to final potential "Ef". However, "Ev" is referred to the 'turn-around' potential for cyclic potentiodynamic scan. The cyclic potentiodynamic curves obtained from as-plated chromium deposits are shown by figures 6a to d vhich demonstrate the results of tests to characterize the performance of chromium deposits in 3* NaCl. The degree of hysteresis indicate the tendency of the plated specimen toward localized corrosion. The hysteresis is due to the accelerated corrosion rate within pits which are formed during the forward scan (13). The values of the electro chemical corrosion parameter for chromium deposits in 3% NaCl with ultrasonic agitation and vithout it are shown in Table 0). From figures 6a - d and Table (3), it can be said that one can deposit a high quality chromium layer with good a corrosion resistance with ultrasonic agitation, especially when it is applied just before chromium deposition. Conclusion 1. Ultrasonic agitation increased the current efficiency and deposition rate for the deposition of chromium from an electroplating bath. The deposit from the plain bath with ultrasound had a smooth surface with a regular and well defined structure. The deposit from the bath exhibited increased microhardness as a result of ultrasound. -290 - Covering Power % 49* 27% 29* 9% 26* 34* Fig (4) covering power using Hull cell under different condition of ultrasonic application. (1) 5 min. ultrasonic (2) 5 min. without ultrasonic (3) 2 min. without ultrasonic +1 min. ultrasonic + 2 min. without . ultrasonic (4) 3 min without ultrasonic + 2 min. with ultrasonic (5) 2 min. without ultrasonic +• 2 min with ultrasonic (6) 1 rain, with ultrasonic + 2 min, without ultrasonic + 2 min. with ultrasonic -291- a. with ultrasonic agitatioa. z 500 b. vithout ultrasonic agitatioa. x 500 Fig. (5): Surface morphology of chromium deposit: a. vith ultrasoaic agitatioa b. vtthout uitrasoaic agitatioa -292- Table (2) Current Efficiency vitb Different conditions of Ultrasonic Agitation Ho. Condition current Effieincy % 1 3 min. Ultrasonic + 2 Bin without 14.27 ultrasonic. 2 2 min. without ultrasonic + 3 ain. 7.28 with ultrasonic 3 1 Din, ultrasonic + 2 Bin. without 11.57 ultraonic + 2 ain. with ultrasonic 4 2 min. without ultrasonic + 2 min. 6.85 with ultrasonic + 1 min. without ultrasonic Table (3) Electrochemical Corrosion Parameters in 3% Had. Ho. Condition Eeorr.V Xcorr. Degree of (SCE) HA/Ca2 hysteresis 1 5 min. Without ultrasonic -0.492 6.8X10* High 2 5 min. with ultrasonic. -0.442 1.20X102 Low 3 2 min. With ultrasonic + -0.456 1.62X.102 low 3 min without ultrasonic. 4 1 min. Without ultrasonic -0.486 2.3X10J High + 3 min with ultrasonic. -293- Potential V Potential V NA/cm NA/cm Potential V Potential V 10 NA/cm2 Fig -6 cyclic potentiodynamic curves in 3% NaCL a- 5 min. without ultrasonic b- 5 min. with ultrasonic c- 2 min. with ultrasonic + 3 min. without ultrasonic d- 1 min. without ultrasonic + 4 min. with ultrasonic -294- 2. The chromium deposits produced by ultrasonic agitation are less porous and are distinguished by higher corrosion resistance properties 3 Ultrasonic agitation tends to refine the grain size of the deposit References 1. T, W. Vie. Y. Y. Wang and C. C. Wan. Plat. Surf, Finish. 66 (3). 47 (1979) 2. A. Roll, Metal Finishing 9. 55<1956> 3. S. Rich. Plating. 42, 1407 (955) 4. R.G. Hickman.Plating52. 407(1965) 5- R. Walker and N. Holt, Plat. Surf. Finish. 68 (2), 44(1981) 6. L. Dominkov, Metal Finishing 20.55 (1967). 7. B Shenoi, Metal Finishing 23.40 (1970) 8. M. S. Frant. J. Eiectrochem. Soc. 108,774 (1961). 9. T. G. Kozan. Plating 49,495 (1962) 10. R. Walker andR. C. Benning, Plating 58. 476 (1971). 11. D. R. Gabe, Principle of Metal Surface Treatment and Protection, Pergman Press (1978). 12. R. Weiner and A. Walmsley. Chromium Plating. Finishing Publications LTD.. (1990). 13. M. L.Rothstein, Plat. Surf. Finish. 73. (11). 44 (1986). -295- £G fSolSMJ First Intematiot. i Spring School & Symposium on Advances in Materials Science (SAMS 94) 15-20 March 1994 Recrystallization of Cold Worked A1-A1203 Alloys And Commercial Purity Aluminium A.S.Taha, S.El-Houte, F.H.Hammad and N.Hansen Atomic Energy Authority, Cairo, Egypt Materials Department Riso National Laboratory, DK-4000, Roskilde, Denmark Abstract Recrystallization and grain growth of commercial purity aluminium and Al-Al20, alloys containing 0.6 and 1.0 wt. * AljO, have been studied after cold-rolling to a reduction of 90*. Heat treatment have been carried out in the temperature range 473 to 893 K and the annealing behavior of deformed specimens has been followed by microhardness testing, optical microscopy and X-ray diffraction studies. It has been observed that the presence of A1,0, particles affects both the kinetics of recrystallization and the recrystallized grain size. It has also been found that the recrystallization temperature significantly can affect the grain size after recrystallization. These observations are analyzed and discussed on the basis of previous studies of the annealing behavior of deformed dispersion strengthened materials. 1.Introduction Deformation and annealing are important processing methods for producing desired properties of materials by controlling their microstructure. For polycrystalline specimens deformed at large strains, such as heavily cold rolled sheet or strip, the dislocation density is very high and a strong orientation or crystallographic texture has developed in the specimen [1] . Upon annealing, a sequence of structural changes will take place, e.g. recovery, recrystallization and grain growth [2]. However, sometimes during the latter process migration is restricted to © minority of boundaries only, so that few grains grow very xarge at the expense of all the rest; this is called abnormal grain growth or secondary recrystallization [3]. A dispersed second phase can have opposing effects on the progress of primary recrystallization of a deformed metal: i) Nucleation of new grains can be accelerated, if the second phase particles are comparatively large [4]. If they are very small, nucleation is retarded or prevented altogether [5]. ii) Growth of new grains is always impeded, because of the drag exerted on migrating grain boundary by dispersed particles, especially if they are small and numerous [6,7]. -297- Zener [8] showed that the pinning force due to the presence of dispersed particles depends on their radius and volume fraction and on the grain boundary energy. This pinning force leads to growth to a limiting grain size which depends on the parameters characterizing the dispersed particles: In recent years computer simulation methods have been developed for studies of recrystallization and grain growth both in single phase materials and in particle containing materials [9-11] . It is not the objective of this paper to discuss the results of such simulations except to note that they have led to a number of predictions which are in good accord with experimental observations of metals containing fine particles. Examples are the relative importance of nucleation rate and growth rate and the magnitude of the limiting grain size as a function of particle parameters. The emphasis on computer simulations in studies of recrystallization and grain growth has lead to a need for experimental data to validate the models. Such results are, however, also necessary in order to further understanding of the complicated processes taking place in real materials. It is this latter objective which has initiated the present study. 2.Experimental procedure A1-A1J0J materials were nanufactured by consolidation of atomized aluminium powder. The purity of aluminium phase in the Al-Al20i materials is about 99.5% and commercial purity aluminium (Al2S) of about the same purity has been included for comparison. The aluminium alloys contains irregular shaped particles of FeAlj of diameter 0.63 urn. The Al-Al20j alloys contain in addition Al20j (in the form of plate or disc shaped particles) . The chemical composition and the structural parameters of the materials are given in Table 1. Table 1. Chemical Composition And Structural Parameters Material AljO, Fe Si p.s. T D wt% wt% wt% urn nm nm AlMD 201 0.6 0.2 0.17 13 15 46 AlMD 105 1.0 0.26 0.18 4.8 8.3 52 2SA1(99.5%) - 0.36 0.16 - p.s. : particle size of powder. T: thickness of Alj03 plates. D: diameter of Alj03 plate. Other impurities: 0.03% max. Cu, 0.02% max. Mn, Mg.Zn and Ti each. The specimens were cold-rolled to 90% reduction and annealed in a muffle furnace at temperatures ranging from 473-893 K for one hour. After annealing, the change of hardness was measured using a microhardness tester (Shimadzu-Japan) with an applied load of 500 g for 10 Seconds. More than 10 readings, were taken and the standard deviation was calculated. The annealed specimens were polished and etched in a solution of: cone. HN05 45 ml, HCl 45 ml, HF (40% sol.) 15 ml, and Feci, 10 g; the solution was diluted to 1/3 its concentration. Microstructural examination was performed with polarized light using an Olympus BH2 -298- metallurgical microscope. Specimens of 99.5* Al were cut in the longitudinal direction while for Al-Al,Oj alloys they were cut in the transversal and longitudinal directions. The grains of the recrystallized 99.5* Al were equiaxed, and the grain size was treasu-ed using linear intercept method. The grain shape of the recrystallized Al-AljO, alloys was frequently cylindrical with the cylinder axis parallel to the longitudinal direction: the cylinder diameter was measured in the transversal section and the cylinder length in the longitudinal section. X-ray diffraction was performed using a Shlmadzu XD-3A diffTactometer with Cu K Nickel filtered radiation. The : scan ranged from 30-BO , with 1 /min scan speed. The specimens were mounted in a special holder. The surface of each specimen was held in a vertical position using a piece of plastic gum. All the samples were marked so as to be examined strictly in the same direction and on the same face to avoid any influence of either, on the obtained results. 3.Results Fig.l shows the change in microhardness with annealing temperature for 99.5 Al,Al-0.fi wt.% AljO, and Al-l.O wt. * A1,0,. The recrystallization temperature is determined from the curves to be 545, 585 and 675 K, respectively. This shows that the recrystallization is retarded in Al-AljO, materials when compared with aluminium, and that the degree of retardation increases with increasing fraction of A1,0, particles. The variation in grain size with annealing temperature of Al and Al-Al,Oi materials is given in Fig.2. For 99.5% Al the change in grain size is small in the temperature ranges 543-723 K and 793- 893 K, but a large increase in grain size takes place when the temperature is increased from 723 to 793 K. For the A1-A1,03 materials the behavior is very different from that of 99.5% Al. For Al-0.6 wt% Al30, the grain size decrease continues „nen the temperature is increased to 793 K and increases sigi-'.ficantly when the temperature is raised from 793 to 893 K. For Al-1 wt* AljOj the effect of temperature on the recrystallized grain size is rather small, however, there is a tendency that a rather limited increase in grain size can be observed when the temperature is raised from 723 to 893 K. Typical micrographs are given in Fig.3-6 for the three alloys. Points to be noted from the micrographs are as follows: 1. The recrystallized grains of 99.5% Al are eguiaxed (Fig.3a) and by raising the annealing temperature to 893 K 1620 C) partial melting is observed (Fig.3d). (This observation of partial melting was unexpected considering the purity of the aluminium). 2. The recrystallized structures of both Al-0.6 wt* A1303 and Al- l.O wt% AljOj show elongated grains with ragged grain boundaries at all temperatures (Fig. 4,5 and 6) .however, with one exception: The heat treatment of Al-0.6 wtV Alj03 at 723 and 793 K results in an equiaxed granular structure (Fig.5-a,b). The elongated grain structure in A1-A1,03 materials has been -299- 70 z 5 60 M M J! so 1 Al M wl •/. Al20j -*—* ! 40 AI-0.6wtV. AljO, o 30 99.5 7. Al 20 -i—« *L 200 400 500 600 700 600 900 Ttmptroturc , K Fig. 1 : Effect of annealing temperature on Ihe mlcrohardneM Ooad 300g> pf' 993% Al Al - Oiwt * A120, and Al-lwt* AlaO, alloyi. -300- 600 700 800 900 Temparaturt, K Fig. 2 :Th§ relation between cjrain size and temperature for 99.5 •/. Al.Al-06 wt % Al203 and Al-1wf/.At a Hoys. -301- *VV^«iH££!L Fig.3: The microstructure of 99.5% Al at different temperatures (rolling plane) . a) 583 K, b) 723 K, c) 793 K, d) 893 K. -3D2- Pig.4: The microstructure of 99.5* Al at different Temperatures (rolling plane). a)583 K b) 623 K c) 673 K -303- -J :v -A* «••' -^ 201 •» Fig. 5: The microstructure of Al-0.6wt% A1,0, at different temperatures (rolling plane). a) 723 K, b) 893 K. 2 00 am i Fig.6: The microBtructure of Al-1 wt% A120, at different temperatures (rolling plane). a) 673 K b) 723 K c) 893 K. -304- attributed [71 to the tendency for the oxide particles to be arranged in a banded structure parallel to the rolling 'and extrusion) direction and to a directionality of the deformation microstructure controlled by the presence of small particles and the deformation mode. X-ray diffraction patterns of the cold worked and other annealed alloys ones are shown in Fig. (7-10). The cold worked Al-Al,0, specimens (Fig.8 and 9) and Al (Fig.7) show similar patterns. The (220) line has a higher intensity than the other lines. However, at 673 K, it is interesting to note that the lines of (220) and (311) almost disappear in Al-0.6 wt% Al20, while only a relative change in line intensities are observed for 99.5% Al and Al-1.0 wt% Al,0, (Fig.10). At this temperature, the two former specimens are recrystallized while the latter specimen is still unrecrystallized. At 723 K the four lines are observed for 99.5% Al (Fig.7) while for Al-AljO, alloys only one line was found. However, at 793 K, the (111) and (200) lines for Al-0.6 wt% Al2Oj (Fig.8) and (311) lines for Al-1.0 wt% AlsO, (Fig.9) have disappeared while still the four lines remain for 99.S% Al (Fig.7). When annealed at 893 K both 99.5% Al and Al-0.6 wt% Al2Oj show a strong (200) line compared with Al-1.0 wt% Al203. It is worthnoting here that all the samples had been marked to examine always the same face az the same direction, before and after annealing. 4.Discussion The reduction of the volume energy of a deformed material is the main driving force of primary recrystallization. The recrystallized volume of a material increases during annealing owing to two processes, the formation and the growth of recrystallization nuclei. The grain size of the structure which forms at the end of recrystallization depends on the ratio between the two rates N and G, the nucleation rate and the growth rate, respectively. Generally for a large ratio between G and N the grain size becomes large whereas a small ratio between G and N results in a small grain size. Both N and G and their ratio depend on a great number of factors [12]. Among these factors are the size and volume fraction of the dispersed particles. These parameters also affect the recrystallization kinetics which can be accelerated or retarded. In general the acceleration is associated with coarse particles (> l urn) and/or wide interparticle spacings [13], whilst retardation is associated with fine particles (< 1 um) and/or small interparticle spacings [14]. In alloys with both small and large particles (like in the case of the Al/AljOj alloys examined), the two types of particles will have opposing effects, with a net result which depends on parameters such as the size and interparticle spacing of the large and the small particles and the degree and mode of deformation [5]. In interpreting our experimental results we must take into consideration that we have in A1-A1,0, alloys two types of particles, coarse particles (FeAl,) and fine particles (A1,0,) . Thus, the presence of large intermetallic particles (FeAlj), surrounded by a deformation zone will lead to a significantly higher nucleation and growth rate near the particles [15] . -305- J 99.5 V. Al o 893 K ;26 10 48 5i 61 72 60 Fig.7: X-ray diffraction patterns for 99.5% Al (cold work 90%) annealed at different temperatures, 723 K, 793 K, 893 K. -306- I Al - 0 6 wf/. Al 0, 793'K o CM CM CO 723 K CW I I ,1 I 29 40 A« 56 66 72 80 .8: X-ray diffraction patterns for Al-0.6wt% Al20j (cold work ) annealed at different temperatures, 723 K, 793 K, and 893 -307- Fig.9 X-ray diffraction patterns for Al-lwt% Al20j (cold work 90%) annealed at different temperatures, 723 K, 793 K and 893 K. -308- 6 23 K Al-lwf/. ALQ 2^3 S CM o> o cs Al_0.6wlV. Al 0 ULJ_I *—s fH_ **^ ^•N SM/ o o 31 1 CM CM *-* CM »^ Al >-^ v» ^ L L 2 9 40 48 56 6 J 72 80 Fig.10: X-ray diffraction patterns for 99.5% Al, Al-0.6wt% Al20j and Al-1 wt% AljOj annealed at 623 K. -309- However, the presence of fine stable particles (Al303) where the mean interparticle spacing is of the order of the subgrain size will retard the coarsening of the subgrain structure, which in turn, inhibits the nucleation of recrystallized grains. The inhibition will increase with increasing Al303 content [16] . The growth of recrystallized grains is in general le^'s affected by the presence of particles on any scale suggesting that retardation of recrystallization is largely a consequence of retarded nucleation [14,17]. This effect of small particles is reflected in the formation of very large recrystallized grains both in A1-0.G wt* Al,03 and in Al- 1 wt* Al203. However, an interesting observation is the decrease in the recrystallized grain size of Al-0.6 wt* AljOj when increasing the recrystallization temperature in the interval 573 to 773 K (see Fig.2). This observation is tentatively explained as a relatively faster increase in nucleation rate than in growth rate when the temperature is increased. The heat treatment at high temperatures (>700-BOOK) has led to the formation of relatively large grains both in 99.5* Al and in Al-0.6 wt* Al,Oj (see Fig.2). However, these observations require further analysis to identify the growth process taking place, for example if it is normal grain growth or secondary recrystallization (abnormal grain growth). Such an analysis requires study of a number of factors which can affect the growth process such as size and volume fraction of Al203 particles, grain size heterogeneity and the crystallographic texture [18-20]. As the growth processes of interest are taking place after the completion of recrystallization it is important that growth studies are carried out on specimens which are heat treated to the end of recrystallization, but not much more. It is therefore necessary to choose and control very accurately both the temperature and the time of the recrystallization heat treatment. A part of such studies is also the formation of the strong recrystallization texture observed both in 99.5* Al and in the Al-0.6 wt* Al203 materials. 5. Conclusions Heat treatment from 473-893 K of 90* cold- rolled Al and Al-Al,03 alloys shows that: 1. The recrystallization process is retarded in Al-AljOj alloys when compared to 99.5* Al. The degree of retardation increases with increasing content of Alj03. 2. The retardation of recrystallization is associated with a marked decrease in nucleation rate rather than in growth rate which is related to a strong pinning effect of A1203 particles on the migration of subgrain boundaries. 3. The retardation of recrystallization has led to the formation of large recrystallized grains in Al- 1 wt* A1,03 independent of the recrystallization temperature. In Al- 0.6 wt* A1203 a more complex pattern is observed with a strong relationship between the recrystallized grain size and the annealing temperature. 4. Further studies will include precise annealing treatments in the high temperature regime to further investigate the evolution of the grain structure and the crystallographic texture duriag growth taking place after the completion of recrystallization. Acknowledgement The authors thank Professor M.El-Sayed Ali, Metallurgy Department, Atomic Energy Authority, Egypt for various help and discussions and Dr. D.Juul Jensen, Materials Department, Riso National Laboratory, Denmark for valuable comments on the manuscript. References 1. H.Hu,"Recovery, Recrystallization and Grain Growth Structures", Text Book, Metals Handbook, vol. 9, Metallography and Microstructures, ASM, 1985. 2. R.W.Cahn,"Recovery and Recrystallization", Physical Metallurgy edited by R.W.Cahn and P.Haasen, Part II, 1595, 1983. 3. G.Gottstien and H.Mecking, "Preferred Orientation In Deformed Metals and Rocks: An Introduction to Modern Texture Analysis", edited by RudeIf Wenk, Academic Press, Inc., 183, 19B5. 4. D.Juul Jensen, N.Hansen and F.J.Humphreys,"Texture Development During Recrystallization of Large Particles", Acta Metall., 33, 2155-62, 1985. 5. A.S.Taha, Ph.D. Thesis, Faculty of Science, Cairo University, Egypt, 1978. 6. N.Hansen and B.Bay, "The Effect of Particle Size on the Recrystallization Temp, of Dispersion Strengthened Aluminium- Aluminium oxide Alloys, ScriptaMet., 8, 1291, 1974. 7. N.Hansen and A.R.Jones, 'Recovery and Recrystallization of Particle Containing Materials", 24eme Colloque de Metallurgie, Institut National Des Science Et Techniques Nucleaires, Saclay, Juin, 1981. 8. C.Zener, See Smith, Trans. Metall Soc. AIME, 175, 47, 1S48. 9. F.J.Humphreys,"Particle Simulated Nucleation of Recrystallization", Recrystallization 90, Int. Conf. on Recrystallization In Metallic Materials, Edited by T. Chandra, A Publication of the Minerals, Metals & Materials Society, 112, 1990. 10. v.Yu Novikov, E.A.Zalem and yu A.Smirnova, "Computer Simulation of Microstructure Evolution During Secondary Recrystallization", Acta Metall. Mater., 40, 3459-64, 1992. 11. P.M.Hazzeldine and R.D.J.Oldershaw,"Computer Simulation of Zener Pinning", Phil. Mag. A., 61, 579-89. 1990. 12. S.S.Gorelik, "Recrystallization In Metals and Alloys", Translated from Russian by A.A.Fanasyev, MIR Publishers Moscow, 1981. 13. P.R.Mould and P.Cotterill, "The Effect of Particle Content and Matrix Grain Size on the Recrystallization of Two-Phase Aluminium-Iron Alloys", J.Mater.Sci., 2,241-55, 1967. 14. P.Cotterill and P.R.Mould, "Recrystallization and Grain Growth in Metals", John Wiley & Sons, New York, 1976. 15. A.D.Rollet, D.J.Srolovitz, M.P.Anderson and R.D.Doherty,"Computer Simulation of Recrystallization-III Influence of Dispersion of Fine Particles", Acta Metall. Mater., 40, 3459, 1992. 16. D.Nobili, F.Mezzetti and E.Susi De Maria, "Study of the -311- Kinetics of Recrystalli2ation of Dispersion Strengthened (Al/Al20j) alloys", J. Mat. Sci., 3, 282-87, 1968. 17. J.l.Brimhall, M.J.Klein and R.A.Huggins, Acta Metall. 14, 459, 19S6. 18. P.A.Beck, M.L.Holzworth and P.P.Sperry, Trans. Metall. Soc. A1ME, 180, 163, 1949. 19. K.Detert,"Recrystallization of Metallic Materials", Edited by F.Haessner, 2nd ed., Reiderer-Verlag Stuttgart, 97,1978. 20. T.Gladman, "Recrystallization and Grain Growth of Multi-Phase and Particle Containing Materials", Edited by Hansen, A.R.Jones and T.Leffers, Riso National Laboratory, Denmark, 183, 1980. -312- EcfSo/iSif First International Spring School & Symposium on Advances in Materials Science (SAMS 94) 15-20 March 1994 EFFECT OF TYPE OF COOLING ON THE PHASE TRANSFORMATION OF ZIRCALOY-2 AND ZIRCONIUM-lwt% NIOBIUM Alia Shafie Taha* 'Associate Prof., Metallurgy Dept, Atomic Energy Authority, Cairo, Egypt. ABSTRACT The effect of type of cooling (air, furnace and water quenching) on themicrostructure in the temperature range 673-1223 K was investigated for Zr-2 and Zr-lwt% Nb. The annealing of the samples were followed by microhardness measurements, optical microscopy and scanning electron microscopy observations. The recrystallization temperature was found to be higher for Zr-2 (823- 873 K) than that for Zr-1 wt* Nb (773- 823K) irrespective of the type of cooling used. Retardation of recrystallization for Zr-2 compared with Zr-1 wt% Nb was considered to be due to the solute content which is higher in Zr-2 than in Zr-1 wt% Nb. Segregation of these solutes along grain boundaries will alter ^crystallization. Phase transformation was observed for both Zr-2 and Zr-1 wt% Nb at temperatures higher than 1000 K. At these temperatures a notable increase of hardness was observed for both alloys. The microstructure changes observed in Zr-2 during heat treatment are due to the precipitation of different compounds of Zr and different alloying elements. The formation of this compounds was accelerated by decreasing rale of cooling. For Zr-1 wt* Nb the presence of Nb accelerates the transformation of martensite and basketweave structure which gave more hardening to die material 1. INTRODUCTION Zirconium alloys have an important application as cladding materials in water cooled nuclear reactors. This is because of their unique combination of excellent corrosion resistance and good strength [1], The most frequently used zirconium alloys are zircaloy-2 and 4 which contain 1.5% tin and minor amounts of iron, chromium and oxygen. In USSR Zr-1 wt% Nb is much used. Tin is reasonably soluble in zirconium and provide solution strengthening up to -313- ~ 1123K. Niobium has only a limited solubility in zirconium and its strengthening effect is due to a dispersion of the precipitates which can be achieved by quenching and aging treatments [2]. The phase diagrams of both zirconium-tin and zirconium-niobium alloys [3] arc shown in Figures 1 and 2, respectively. In zirconium alloys, microstructural control by thermal mechanical processing is of particular importance since the selection of alloy composition is severly restricted by neutron absorption considerations. At temperatures lower than 1073K, u.c changes in the microstostructure are associated with the removal of cold work (i.e. recovery and recrystallization) [4.5]. At temperatures higher than 1073K, the changes in the microstructure are associated with any phase changes [5.6]. The cladding material is commonly used in the cold worked condition and could experience microstructural and property change during temperature transients experienced during such occurrences as a postulated loss-of coolant accident (LOCA). It is the aim of this paper to follow such microstructure properly changes using microhardncss, light microcopy and scanning electron microscopy. 2. EXPERIMENTAL The zircaloy-2 and zirconium-1 wt% niobium ingots used in this investigation were formed into rods. The rods were hot and cold rolled to produce sheets. The sheets were annealed for 30 minutes at 973k in vacuum. The resulting average grain size for Zr-2 and Zr-1 wt% Nb was found to be 13 um and 10 |im, respectively. The chemical compostion of Zircaloy-2 and zirconium-niobium alloys is shown in table 1. Table 1 : Chemical analysis of Zirconium alloys in weight%. '— — - Alloying element Material Nb ft Ni Q Qi Tl Al Kb Pb Sn Zircaloy-2 0.13 0.05 0.12 liO Zirconium- 1.0+0.1 0.023 0.019 0.006 0.004 0.003 0.003 0.003 niooium ^ .... -314- Atom* Percent Tin 10 20 ]0 <0 SO SO 70 SO 90 100 "IKS 1171 2071 LUSTS... • 1071 •1673 igpc •1471 -1171 • 107] ,_ • S71 \ .to •npc"?' to so to 100 Sit Weight Percent Tin Fiat: Phase diagram of Zr-Tin alloy. Atomic Percent Niobium 30 U) SO SO 70 SO 90 100 isoo 1771 1100 (-3573 1100 1171 _ isoo 1-1171 £ 1700 I97J 5 1500 1771 Mt,tHb) ! 1300 J-1S73 E 1100' [-1371 "" 900 1173 700 173 10 20 10 tO 50 SO 70 (0 90 100 Weioru Percent Mobkn Nb Figi Phase diagram oTZr-Nb alloy. -315- The specimens were isochronal annealed in a special designed oven; its temperature was automatically controlled to an accuracy ±5'C, at temperatures ranging from 673 to 1223K for one hour in vacuum 5.I0"4 mm Hg. They were subjected to different cooling rates (air-cooled, fumce-cooled t water-quenched). Shimadzu microhardness tester model (Japan-Vicker*s hardness) was used in this investigation. An applied load of 300g for 10 seconds was used. More than 10 reading were taken for each specimen and the standard deviation was calculated for every temerauire. To examine the microstructure by metallogaphy or scanning electron microscopy, the specimena were polished and then etched in solutions of 10 parts HF, 45 parts NH03and45 parts distilled water. After the specimens were etched, they were dried and examined microscopically using Zeiss optical microscopy and scanning electron microscopy (JEOL-Iapan). 3. RESULTS 3.1. Microhardness Measurements The change of hardness with temperature for Zr-2 and Zr-1 wt% Nb followed bu different cooling rales (air-cooled, furnace-cooled, and water-quenched) is show in Figs. 3 and 4, respectively. Figure 3 indicates that for Zr-2 no change in hardness was observed with temperature up to 1023K for furnace-cooled and water-quenched specimens while a slight decrease was noted for the air-cooled specimens. At temperatures higher than 1023K an increase of hardness was observed reaching values of 715 ± 35.0 VHN and 756 ± 85.0 VHN for furnace-cooled specimens at temperatures 1173 K and 1223K, respectively. For Zr-1 wl% Nb (Fig. 4), a slight decrease of hardness up lo 923K followed by increase was observed for all types of cooling used. The main features noted from the microhardness study were as follows: (a) Decrease of hardness for both Zr-2 (673-1023K) and Zr-1 wt% Nb (623-923K) was observed irrespective of die type of cooling used. Complete recrystallization was achieved for Zr-2 at 873K while it was achieved at 823K for Zr-1 wt% Nb i.e. presoftening before recrystallization was occurred and this is in agreement with other results (7). (b) At temperatures higher than 1023K, for Zr-2, the hardness values were frequently higher for furnace-cooled than for air-cooled and water-quenched. However, for Zr-1 wt% Nb, at temperatures higher than 923K, the hardness values were higher for water-quenched than that -316- of air-cooled and furnace-cooled ones, (c) The hardness values of Zr-1 wt% Nb alloy were higher than that of Zr-2 alloy at all temperatures and types of cooling except at 1173K and 1223K furnace-cooled. 3.2. Optical Microscopy Observations Presentative micrographs are given in Fig. 5-7 for Zr-2 and Figs. 8-11 for Zr-1 wt% Nb. Points to be noted from the micrographs are as follows: 1. The microstructure of Zr-2 annealed at 1123K and air-cooled consists of a and transformed P (Fig. Sa), while the furnace-cooled ones (Fig. 5b) consists of a structure (mixed structure of basketweave and parallel plate Widmanstfltten structure). Irregular structure was observed for Zr-2 annealed at 1173K air-cooled and furnace-cooled specimens (Fig. 6). Annealing at 1223K, the air-cooled specimesn showed distorted a elongated grains (Fig. 7a), while the furnace-cooled specimens showed equiaxed a-phase and some transformed P-phase (Fig. 7b). 2. For Zr-1 wet% Nb annealed 1123K and air-cooled, the microstructure consists of irregular structure (Fig. 8). Annealing at 1173K, brought it into the a + fi-Zr phase region and the P-Zr was higher than after treatment at 1123K. The P-Zr is unstable and transormed, at least partially, to an a structure on air cooling (Fig. 9 a, b and c) or to martensite structure on furnace cooling (Fig. 10 a, b and c). Annealing at 1223K (single-phase P-Zr region) the air-cooled specimens showed basketweave structure (Fig. 1 la), while the furnace-cooled ones showed parallel-plate WidmansUUten structure (Fig. lib). 4. DISCUSSION It appears from the results that, die changes in microhardness for Zr-2 and Zr-1 wt%Nbwiih annealing in the temperature range 673-1223K and followed by air, furnace cooling and water quenching, can be divided into two categories. The first one in the temperature range 673-1023K and is related to recovery and recrystallization (rearrangement and annihilation of dislocations) while the second one in the temperature range 1023-1223K and is related to any phase changes. Microhardness measurements and optical microscopy were not able to reveal the presence of recovery which involved the rearrangement of dislocations. Thermoelectric -317- I I I • Air.coolid o Furnoce - cooled » Woltr- quenched Zr-2 ^tl-tt# aoo soo 1000110 0 1200 Temperolure. K Fig 3: The variation of micf ohardntu (load 3O0g) with temperature tor Zr-2. -318- Ol I ' ' 600 700 (00 900 WOO ttOO 1200 Ttmptratur*, K Ffc.4: Tha variation of mfcrohardnast (load 300j) with lampantura (or Zr-1 wlKNb. -319- fia. 5 : The microstnictiiic of Zr-2 annealed al 1123K for one hour, a) air cooled, and b) furnace-cooled. Kg. 6 : Tneinictostn)C&in:ofZr-2annea) -320- Fig. 7 : The microstructure of Zr-2 annealed at 1223K for one hour, a) air cooled, and b) furnace-cooled. Fig. 8 : The microstructurc of Zr-1 wt% Nb annealed at 1123K for one hour and air-cooled. -321- Fig. 9 : The microstructure of Zr-1 wt% Nb annealed at 1173K for one hour and air-cooled. a) optical microscopy, b) and c) scanning electron microscopy (different places). -322- power and electrical resistivity measurements, were able to monitor these changed during recovery (4,7,8). In this work, complete recrystallization was achieved at temperatures 873K and 823K for Zr-2 and Zr-1 wt% Nb, respectively. This indicates that, the recrystallization temperature is slightly increased over that of pure Zr (9,10). Increase in rccryslallizalion temperature of Zr has been observed with the addition of any of the following solutes Ti, V, Cr, Mn, Fe, Ni, Mo, Hf, Nb and Sn (1). So the retardation of recrystallization in Zr-2 than Zr-1 wt% Nb could be attributed to the solute content which is higher in Zr-2 than Zr-1 wt% Nb. During annealing, solutes lend to segregate at interfaces(ll) such as grain boundaries which are preferred sites for nuclcauon thus alter recryslallization. Nonhwood et al (5) attributed die large decrease of microhardness in the temperature range 673-873K for Zr-1 wt% Nb to the beginning and complete crecrystallizalion of this alloy. The microstructure investigations for both alloys at temperature in Ihe range 1123-1223K indicates that, there are two types of transformation can take place which depend on Ihe composition and type of cooling. The first one was found al temperature 1123K (o + B phase) which is just below the B-phase while die second one at temperatures higher than 1123K which is in the B phase. Higher values of hardness for both alloys were obtained if they were solution treated in the B-phase region rather than in the a + B phase region. This is inagreement with previous work on Zr alloys (12-14). For Zr-2, it was found that Ihe a-region extended up to 1053-1083K, the B region from 1253-1258K to high temperature, and that intermetallic precipitates were formed at 11S3K (6). The particles are finely dispersed in the matrix, and their sizes are at most a micron (IS). Thus the formation of the intermetallic precipitates at temperature higher than 1123K gave strengthening to the material. For Zr-1 wt% Nb, it was observed from optical microscopy study that the volume fraction of B-Zr was higher after treatment at 1173K than at 973K (S). Intermetallic precipitates were formed at ct-Zr grain boundaries along with the B-Nb in annealed materials (16). Selected-area electron diffraction identified these precipitates as Ihe Zr (Fe, Cr), laves phase type (17). Both Fe and Cr are b-stabilizers (14) and would tend to concentrate in die B-phasc where they would have a greater solubility than in ot-Zr phase. On annealing, the volume fraction of the B-phase decreased, and the Nb content in die B-phase increased as B-Zr transformed to B-Nb. These changes made it more difficult for the Fe and Cr to remain in solid solubility in the b-phasc and they were precipitated as die Zr (Fe, Cr). Laves phase (16). Also the P-Zr is Fig. 10 : The microstructure of Zr-1 wt% Nb annealed at 1173K for one hour and furnace-cooled. a) optical microscopy, b) and c) scanning electron microscopy (different places). -324- H1 ...(«•" Fig. 11: The microstructure of Zr-1 wt% Nb annealed at 1223K for one hour, a) air-cooled, and b) furnace-cooled. -325- unstable and transforms (5), partially, to an a structure (basketweave or parallel plate Widmanstauen structure) on cooling at RT. It was found before (12) that samples have a basketweave a-structure, with the possibility of a small amount of retained p-phasc, or the brittle h.c.p. w-phase (14), had higher strengths but lower ductilities. This is in agreement with our results for Zr-2 and Zr-1 wt% Nb. 5. CONCLUSIONS Annealing cold worked Zr-2 and Zr-lwt% Nb sheets in the temperature range 673-1223K followed by different types of cooling, two main categories of change in the microstructurc can be distinguished: 1. The first one in the temperature range 673-1023K is associated with the removal of cold work, i.e. recovery and rccrystallizalion. Retardation of rccrystallizalion for Zr-2 compared with Zr-1 wt% Nb was atributed to segregation of solute atoms which are higher in Zr-2 along grain boundaries. 2. The second one in the temperature range 1123-1223K is associated with phase changes. Higher values of hardness for both alloys were obtained if they were solution treated in theP-phase(>1123K). 3. Annealing at temperatures higher than 1123K accelerate the formation of precipitation of intcrmetallic compounds or basketweave and martensetic structure which give strengthening to the material. REFERENCES 1. D.L. Douglass, The Metallurgy of Zirconium, Atomic Energy Review Supplement, IAEA. Vienna. (1971). 2. R. Krishnan and M.K. Asundt, Zr Alloys In Nuclear Technology, Proc. Indian Acad. Sci., (Engy. Sci.). 4(1) (1981) 41. 1 B.M. Thaddcus, L.M. Joanne and H. Baker, Binary Alloy Phase Diagrams, 2 (1987) 2087. 4 P Merle and R. Barrelly, On The Possibility Of Studying Recovery And Recrystallization Phenomena By Thermoelectric Power Measurements, case of Zirconium Alloys. Scripta Mctallurgica, 20 (1986) 1089. -326- 5. D.O. Northwood, J.W. Robinson and Z. Jic, Recovery, Rccrystallizalion And Phase Transformations During The Annealing or A cold-Worked Zirconium Alloy, Rccryst. 90, Edited by T. Chundra, The Minerals, Metals & Materials Society, (1990) 28S. 6. D. Arias and R.C. Gucrra, Phase Transition Temperature In Zircaloy-2, J. Nucl. Mat., 144 (1987) 196. 7. OJL. Gray, Recovery And Recrystallization of Zirconium And Its Alloys-Part II: Annealing of Cold-Worked Zirconium, Report HW-69679 (1961). 8. P. Mexle, C. Vauglin, G. Fantozzi, J.L. Derep and D. Charquet, Study By Thermoelectric Power Measurements of Recovery and Recrystallization of Cold-Rolled Zr-4 Sheets-Zirconium in the Nuclear Industry, Seventh Internation Symposium, ASTM STP 939, R.B. Adamson and LJF.P. Van Swan, Eds. American Society For Testing And Materials, Philadelphia, (1987) 555. 9. J.C. Goodwin and K.M. Goldman, Production annealing of Zr-2, WAPD-120 (19SS). 10. S.H. Bush and R.S. Kemper, Recovery And Recryst of Zirconium And Its Alloys-Part 3 Annealing Effects In Zr-2 And Zr-3. HW-69680 (1961). 11. R.W.K. Honey Combe. The Plastic Deformation of Metals, Edward Arnold Publishers, Ltd, London, (1968) P. 80-100, and 299. 12. D.O. Northwood and M. Thit, Microstructure-Mechanical Property Relationships In A High Stength Zr Alloys, Microstructural Science, 9 (1981) 277. 13. D.O. Northwood, Heat Treatment, Transformation Reactions And Mechanical Properties of Two High Strength Zr Alloys, J. of the less-Common Metals, 61 (1978) 199. 14. D.O. Northwood and D.T.H. Lim, Phase Trans. In Zr And Its Alloys, Canadian Metallurgical Quarterly, 18 (1980) 441. 15. T. Rubo, Y. Wakshima, H. Imahashi and M. Nagai, Distribution of Intermetallic Particles And Its Effects on SCC of Zr Alloys, J. Nucl. Mater., 138 (1986) 756. 16. PJ. Kenny, G.S. Cole, D.O. Northwood, J. Wylie. and G.F. Vander Voort. A Metaltographic Study of the Mkrostructure of a Zr-2.5 wt% Nb alloy. Microstructural Science. 17(1989)459. 17. X-Meng and D.O. Worthwood, Polytype structures In Zr-Cr-Fe Laves Phase, J. Less Common Metals, 125 (1986) 33. -327- EC fSo^o First International Spring School & Symposium in Advances in Material Science (SAMS 94) Cairo, 15-20 March 1994 MICROSTRUCTURAL CHARACTERISTICS OF SPLAT COOLED Al-Cu ALLOYS A Abbas, R.M. Ramadan, AE. El-Nakhali, and S. Ibrahim Faculty of Pet. & Min. Eng^ Suez Canal University, Suez, Egypt. Abstract Microstructural characteristics of splat cooled Al-Cu alloys, 5 to 30 wt% Cu have been studied. Splat cooling was carried out using the twin roll technique. It was generally noticed that the microstructural features were almost uniform in the cross section irrespective of the ribbon thickness, which was in the range 30 to 70 um. A single phase structure with thick boundaries of solute atoms segregation was observed at a close range of thickness 40 to 60 um of specimens containing up to 24% Cu. Eutectic, was observed locally in the thinner areas of the 15 and 24% Cu specimens. For both 5 and 15% Cu alloys, fine Al2Cu precipitates at grain boundaries, in the thicker areas were noticed. Eutectic become dominant for the 27 and 30% Cu. At 27% Cu, degenerate eutectic and small amounts of radial lamellar or radial rosette structures were observed, whereas, at 30% Cu the radial rosette morphology was predominanting. Finally, the produced splat thickness and the observed microstructural features indicated the achievement of a very high rate of cooling rates of about 10^ to 109 K/s using the twin roll technique. 1. Introduction The aluminum-copper system, in particular the eutectic composition, has been extensively studied in the as-splat condition'1"^. However, more effort is needed to have a -329- global idea of the hypereutectie zone structures during rapid solidification. There is still considerable uncertainty concerning the microstructure characteristics of these alloys, particularly with certain morphologies and the distribution of the phases and their relation to cooling rates. The techniques of splat-quenching of metals and alloys directly from the liquid state to room temperature or below, are well established'4'. However, the twin-roll technique is less commonly used in the splat cooling process. There are reasons to believe that this technique is more effective for rapidly cooling the specimens'5^. The purpose of this work is to examine the effects of copper content changes in the Al-Cu system on the structures morphology when using the twin-roll splat technique. The results are discussed in the frame-work of the existing solidification model and previous results on the eutectic composition, also the effeciency of the twin roll technique is to be estimated. 2. Experimental Work Nine alloys, from 5 up to 30 wt% Cu(5,9,12,15,18,21,24,27, and 30 wt% Cu) were prepared from commercial pure aluminum and 50/50 Al/Cu master alloy in an electrical resistance furnace. The chilli cast alloys were splat quenched from temperature 150 C above the liquidus for all compositions. Details of the twin - roll apparatus are given elsewhere (6). The splat strips did not adhere to th copper substrate durng the splat process. The otained specimens are extremely brittle at the high copper content and become almost ductile for the lower copper concentrations. The specimens were mounted for both microstructure observation and microhardness measurements. -330- 3. Results & Discussion The Homogeneity of the As-splat Products The splat cooled specimens obtained using the present technique in the from of ribbons, varied from 5 to 12 cm in length and about 2 to 3 cm in width. A variation of splat thickness from 30 to 70 um, was obtained which indicates that the cooling rate was in the order of 106 to 109 k/min. The estimation of the cooling rate was based on the thickccs evaluation method by Jones' . However, the thickness could not be fully trusted as an accurate estimate of cooling rate, since the cooling efficiency is determined by thermal contact and consequently by the heat transfer coeficient between the splat and the substrate which is technique dependent. Furthermore, it was generally observed that the liquid metal viscosity was decreased as the copper content was increased which would contribute to the observed thickness variation. It was generally noticed that the microstructure in the cross section was almost uniform, irrespective of the ribbon thickness, which indicates a homogeneous cooling rate all along the cross section. This is one of the advantages of the present technique. Nevertheless, a deviation of such homogenity has been observed for the near eutectic composition, between 24 up to 30 wt % Cu- This deviation could be due to a high level of solute content, which is likely to undermine the degree of undercooling and consequently the cooling rate in general. It is worth noting that, in using other splat techniques, the structure homogenity was not obtained'8"10'. Variation of Microstructure with Composition : With such a wide range of compositions, 5 to 30% Cu, an extended range of phases and morphology would be expected. However, observation indicated a limited number of microstructural features. A dominant supersaturated single phase and/or radial lamellar, radial rosette or -331- degenerate eutectic are generally observed depending on copper concentration. However, it is possible to observe from Fig 1 that the so called single phase, appears to contain thick boundaries. Such boundaries are not likely to be of a second phase since the presence of such a phase has not been reported previously under either similar rate of cooling or specimen thickness. Therefore, it is possible to suggest that these thick boundaries are due to solute atoms segregation, which may take place during or after solidification. Similar observation has been made^ ' and confirmed the presence of solute atoms segregation close to grain boundary areas of Al-4.5% Cu alloy at a cooling rate of = 105 K/s. In general, such a single phase structure is mainly observed at a close range of thickness, 40-60 (am, for specimens with different copper contents up to 24% Cu. Moreover, an eutectic structure morphology was observed locally in the thinner area of the cross section for specimens having copper content between 15 and 24% Cu. This could be due to local solute segregation in the liquid. The general microstructural characteristics resemble those previously reported for 5-12 at % Cu (10-24wt% Cu) alloyst2]. The effects of the rapid solidification structure on the stength of Al-Cu alloys are shown in Fig. 2 as a function of the copper content, the microhardness results are compared with those of the normal chill cast alloys. The relatively higher hardness values of splat specimens compared to those of the chill cast ones, could be due to several possible strengthening mechanisms, such as solute solid solution strengthening, structural refinement and high concentration of point and lines defects due to the high rate of cooling. Specific features were observed for both 5 tand 15 wt% Cu splat alloys, where grain boundary precipitations frequently occured in the thicker areas as shown in Fig. 3a and b. The precipitation appears to be very small and most likely of the equilibrium q-phase (Al2Cu). similar observations were made for the alloys of 4.5 wt% Cu'11,12' -332- && x Fig. (1) : Typical of the as-splat cooled microstructure on the specimen surface, 15% Cu, (1250X). 400 a as-chill cast • as-splat cooled 300 I 200 n 10 V c 15 U 3ID 100 0 20 30 wt.% Cu Fig. (2) : Microhardness of both the as-chill cast and the as-splat cooled alloys with the weight percent of copper. -333- and 15% Cu'13'. However, it is not likely that the the presence of grain boundary precipitates is related to such particular alloys compositions. A reasonable explanation for the appearance of this type of precipitates at such high cooling rate was proposed by Williams and Edington'13'. Using a point to point analysis on the foil, they indicated that there was no trace of diffusion which might have occured in the solid states, and the solute rejection occured into the liquid phase during solidification leading to the nucleation of 0-phase. Therefore, it is possible to consider that the nucleation of 9-phase could rather occur in the liquid state, as will be pointed out later while discussing the nucleation of the degenerated eutectic structure. It is important to point out that grain boundaries precipitates were not observed in thinner areas of both 5 and 15 wt% Cu aloys, as well as other alloys compositions irrespective of their thickness. It is therefore possible to speculate that the reason behind such phenomena could be related to one of cooling rate parameters, and/or a possible relatively prolonged holding time prior to solidification. The Eutectic Microstructure: As indicated before, the eutectic morphologies were observed only occasionally in compositions above 15 wt% Cu, and became dominant for the 27 and 30 wt% Cu. A degenerated eutectic type of structure was observed in the splat cooled 27% Cu alloy in addition to a relatively small amount of radial lamellar or radial rosette structures, as showin in Fig. 4a. Whereas, in the 30 wt% Cu the commonly observed structure is rather the radial rosette morphology and a minute quantity of both radial lamellar and degenerate structures, Fig. 4b. Similar structure features have been observed for the eutectic composition, 33 wt% Cu(l,2,14] The fact that both lamellar and radial rosette types of structures appeared on both cross section and foil surface -334- Fig. (3) : Grain boundary precipitation o-id the thicker zones for both a) 5% Cu (800X), and b) 15% Cu (800X). -335- Fig. (4) : Radial lammella and radial rosette structure for both a) 27% Cu (1125X), on the specimen cross-section. b) 30% Cu (1000X), on the specimen surface. -336- suggested that they assumed a nearly spherical arrangement in the material. Analyses which have been earned out'13', for the particle at the centre of the rosette structure indicated that it was supersaturated copper in aluminum, with about 5 at %Cu. This idnicates once more that solid state diffusion does not occur during solidification in order to reach the equilibrium condition. Furthermore, it would also suggest that the nucleation of such phase morphology occurs in the liquid. On the other hand, a similar line of conclusion could be drawn in what concerns the nucleation process of the degenerate eutectic structure. However, it should be stressed that both lamellar or rosette structures and degenerate eutectic form at different rates of cooling. The latter is reported to occur at a higher rate of cooling^2'111. Moreover, the present results are in agreement with the previous^ work indicating that the eutectio coupled zone in the Al-Cu system lies in the range of 24.5-40 wt% Cu^, for the high rate of cooling, therefore, confirming the high cooling rate capabilities of the twin roll technique used in the present work. 4. Conclusions Based on the obtained results and the reported data it is possible concluded that: 1. A cooling rates in the range of 106-109 K/min was achieved using the twin roll technique. 2. High hardness values of splat cooled materials were observed to be due to several strenthening mechanisms, e.g. extension solid solubility and fine dispersoids. 3. The eutectic morphology detected on various copper concentraions indicated the extension of the eutectic coupled zone in the Al-Cu system. -337- 5. References /. M.H. Burden and H. JOns, J. Inst. Metal, 98(1970) 249. 2. D.B. Wiliama and J.W. Edington, J. Mater. Sci.,12 (1977) 126. 3. M.G. Scott, ibid, 9 (1972) 1372. 4. T.R. Anatharaman and C. Suryanaryana, ibid, 6 (19710 1111. 5. K. Miyazawa and J. Szekely, Met. ???, 124 (1981) 1047. 6. A. Abas, M. Sc. Thesis, Suez Canal University, 1990. 7. H. Jones, Rapidly Solidification of Metals and Alloys, Inst, of Met., London, (1982), 1-14. 8. A. Munitz, Met. Trans., 16B (1985) 149. 9. B.A. Mueller, J.H. Perepezko, J. Mater. Sci. Eng., 98 (1983) 153. 10. S.C. Huang and H.C. Fiedler, Met. Trans, 12A (1981) 1107. 11. M.G. Scott and J.A Leake, Acta Met., 23 (1975) 503. 12. L.J. Masur and M.C. Flewings, Proc. 4th Conf on RQM, (1981) 557. 13. D.B. Williams and J.W. Edington, Proc. of 2nd Conf. on RQM, (1976) 135. 14. M.G. Scott, J. Mater. Sci., 10 (1975) 269. -338- First International Spring School & Symposium on Advances in Materials Science (SAMS 94) 15-20 March 1994 FORMATION AND THERMAL STABILITY OJ; ICOSAHEDRAL PHASE IN RAPIDLY SOLIDIFIED Al-15Mn-2Ti ALLOY R.M. RAMADAN FACULTY OF PETRO.&MW. ENC, SUEZ CANAL UNIV., SUEZ, EGYPT. ABSTRACT: The formation and thermal stability of the icosahedral phase (i-phase) in ternary alloy, AM 5 wt%Mn-2 wt%Ti produced by melt spinning technique are investigated. In order to monitor the thermal decomposition of the melt spun material several techniques are employed. X-ray diffraction (XRD), differential scanning calorimeter (OSC), micro-hardness, and optical microscope (OM). The obtained results show that the as-melt spun product contains only the i- phase embedded in a matrix of super saturated solid solution (SSSS) of Al. During annealing the SSSS firstly decomposes into metastable phases, G-phase and metastable AlgMn. Next, the i- phase decomposes into metastable AI6Mn and another metastable phase called t-phase. Finally, all the metastable phases transformed into the stable AI6Mn. it is concluded that the addition of Ti thermally stabilizes the i-phase as well as influencing its thermal decomposition behaviour. INTRODUCTION: The icosahedral phase (i-phase) has been reported to form in the rapidly solidified Al-alloyed with Mn, Fe, Cr or V[J-9). This phase exhibits an electron diffraction pattern which has sharp spots, but has five symmetry axes which are not consistent with crystal lattice transitions. During rapid solidification the i-phase grows dendriticafty and competes with several other stable and metastable phases for dominance in the microstructure (depending upon the percentage of the transition metals in the alloy and the applied cooling rate). The icosahedral phase was observed in some regions of the ribbons with as little as 5 at* Mri[3,4] and it increases in abundance as the Mn content increases[5]. A few studies have reported the thermal stability of the i-phase as a function of alloy composition for the range of Mn contents up to 22 at%[5,6]. The present work reports the influence of the presence of a third element, titanium, on the formation and thermal stability of the i-phase formed in previously investigated AM5wt%Mn{7-Q] with the phases with which it competes. Titanium was chosen because it forms stable and metastable phases with both AJ and Mn (AI,Ti and Mn9Ti) and has a low diffusivity coefficient in the Al lattice which is anticipated to thermally stabilize the as-melt spun structure. -339- EXPEREVIEi\TAL PROCEDURES: An alloy of nominal composition AI-15wt%Mn-2wt1tTi was prepared using commercially pure Al and the commercial master alloys AI-5wt%Ti and Al-75wt% Mn. An ingot of the required alloy was made by casting from maximum of 1000 *C into a steel mould. The ingot was further homogenized at 625 *C for 72 hr and sectioned for a melt spinning process. The melt spinning was done in a chamber under He atmosphere. Each sample was remelted in a quartz tube heated by a high frequency induction current and the melt stream was ejected by He gas onto a copper wheel, 140 mm in diameter and running at a rotating speed of 3500 rpm. Ribbons of 30- 70 um in thickness and 6-10 mm in width were obtained. The study was carried out using thin ribbons in the range of 30-45 pm. Thin ribbons were cut into small lengths of 30 mm and annealed in a vacuum quartz tube at temperatures of 350, 400, 450, 550, and 620 *C for several intervals of time. The techniques used to investigate the thermal stability of the melt spun ribbons were: X-ray diffraction (XRD), differential scanning calorimetry (DSC}, microhardness measurements, and optical microscopy. X-ray measurements on both as-melt quenched and annealed ribbons was carried out at room temperature using Cu Ka, radiation. For OSC, samples were heated at rate of 10*C/min in pure argon up to 700 *C. Microhardness measurements were carried out using a load of 25 gm on both as-melt quenched and annealed ribbons. The microhardness value was taken as an average of 20 readings. Microstructures of as melt-quenched as well as annealed ribbons were investigated using optical microscopy Nephot 21 on a cross-sectioned surfaces etched by Killer's reagent RESULTS AND DISCUSSIONS: X-ray Data: The X-ray diffraction data of the as-melt quenched ribbons and the annealed ones are given in Tables l-lll. The X-ray patterns of both as-melt quenched and annealed ones are shown in Figs 1 and 2. The different phases formed in the investigated ribbons were estimated from the measurements of diffracted X-ray lines using scanning range of 10< 26 <115. Because samples were not pulverized results were somewhat influenced by preferred orientation textures which may result in some discrepancies in the respective intensities of the detected phases. The equilibrium limit of solid solubility of Mn in Al is approximately 1.6wt%, but this solubility extended by melt quenching technology[10]. The manganese content of the a-AI present in melt quenched ribbons can be deduced from the lattice parameter measurements of a-AI. Figure 3 shows the measured d-spacing of the (331) plane of Al present in the as-melt quenched ribbons as well as annealed ones. A comparison with the results obtained by Ohashi et alr11] (that illustrate the d^ of the Al-lattice as a function of the manganese concentration) shows that the present measured value of the d^-j of Al-lattice in the as-metal quenched ribbons is indicative of a manganese content of about 5wt% Thus in this alloy the i-phase must be growing with a manganese content greater than the average composition of the alloy leaving the intercrystalline liquid depleted in manganese. The interplanar spacing d331 is observed to increase significantly with the annealing temperature owing to the precipitation of the manganese content when annealing temperatures exceed 350*C. The measurements of d33i values, Fig 3, displays a complete decomposition of the a-AI after annealing at 400*C for two hours. -340- Phase identifications: The X-ray data obtained for the as melt-quenched AI-28wt%Mn{2], AI-15 wt % Mnf7], and Al-(18- 25) wt% Mn[12] alloys are given in Table I for comparison. As can be seen, (except the first line corresponding to d-spacing of 0.2314 nm[7]) a similar pattern is observed in the present case. This line has not been reported by others(2,3,12J. The obtained X-ray data reveal the presence of the quasicrystalline phase, which is called i-phase embedded in a matrix of super saturated solid solution of aluminum Figure 4 shows the microstructurc of the as melt-quenched material Tables il and III and Figs 1 and 2 demonstrate the X-ray results obtained for annealed ribbons At 3S0*C and for times as little as one hour annealing does no! yield a drastic change either in the position of the i-phase lines or in their relative intensities. However, reflections of the metastable 6-phase and the metastable AI6Mn are detected. The strong lines of these reflections come from the planes (310), (222), and (321) of 6-phase at d-spacing of 0.2372, 0,2155, and 0.2008 nm respectively after annealing at 450*C for one hour. Therefore, it can be deduced that the formation of the metastable phases at such an annealing temperature results from the decomposition of the a-AI SSSS rather than the crystallization of i-phase. Increasing the annealing temperature up to 450*C or the annealing time at 400°C for up to four hours results in the diffracted lines of the metastable phases ( G-phase and AlsMn ) appearently gaining in intensity, (see Tables II and III). At this stage of annealing the formation of the metastable Al6Mn is observed to associate with the elimination of the i-phase. It is interesting to observe that annealing for one hour at 450*C is found to be unsufficient to eliminate completely the reflections of i-phase, Table II. Such results are utterly in disagreement with those obtained from the binary alloy AI-15wt%Mn[7] which show complete elimination of the i-phase after one hour annealing at 400*C . Based on their X-ray data of melt-quenched AI-Mn-1wt%Zr containing up to 13 wt% Mn alloys Ohashi et al, have speculated that a precipitate forming during the early stage annealing of the melt-quenched ribbons up to 400*C is a metastable phase or a pseudoamorphous phase of the equilibrium AlcMn because their X-ray diffraction lines could not be assigned to any other metastable phases, eg C, G, and G" phases. Thereto, they anticipated that the speculated phase has a lattice sue somewhat larger than the equilibrium AI6Mn. It is the metastable AI6Mn phase which is reported to form in thin films of Al- Mn alloys with an orthohombic unit cell (a=1.525, b=1.240, and c=1.783nm)[13] compared with the orthohombic unit cell of the stable AI£Mn (a=0.6S, b=0 775, and c=0 887 nm)[l4]. As a result of the differences in the unit cell dimensions between the two mentioned phases it can be assumed that some coherency strains are associated with the precipitation of this phase and that these facilitate the hardening observed in the microhardness values of the heat treated ribbons containing metastable AI6Mn, Fig. 5. In addition to the metastable AIGMn, the i-phase is observed to decompose to another quasicrystalline metastable phase i.e. t-phase. Such process occurs after annealing at 400*C for four hours or at 550*C for 30 min ( see Tables II and III) which could will be confirmed through X-ray analysis. However, only four lines out of decagonal t-phase could be detected and they have the following d-spacing: 0.2246, 0.2089, 0.2072, and 0.2051 nm[3]. Furthermore, among those four lines of reflection, two lines ( 0.2089 and 0.2051 nm ) are overlapped with the lines of other competitive phases, stable AI6Mn and G-phase. The other two lines normally detected without overlappmg[3]. The presence results confirmed the presence of those two lines at d- spcing of 0.2246 and 0.2072 nm. Therefore, it is possible to confirm the existance of the t-phase. The thermal analysis results support this conclusion as will be seen, later. The decagonal phase is previously detected in an annealed melt-quenched AI-5wt%Mn ribbons[8] and latter is observed in the as melt-quenclrjo' Al-Mn ribbons containing 30 or more wt%Mn and is found to be especially prominent in the thick ribbons[3 and 5], Decagonal phase is observed to grow expitaxially on the i-phase[3J. The X-ray data show that the t-phase disappears after annealing at 550*C for 24 hours or at 620°C for one hour, Table III. This results are in a good agreement with -341- the DSC data, Fig 6, as will be discussed ater. At this stage of prolonged annealing at S50*C (for 24 hrs) or high temperature annealing at 620"C for one hour, the detected phases are a-AI and the equilibrium phases AlsMn, Mn3Ti, and AI3Ti. Reflected *'nes from planes of AI,T< are weak and few as a result of its small volume fraction. Thermal Analysis: The thermal analysis of the as-melt quenched ribbons was camedout using the DSC technique A typical DSC curve at heating rate of 10*C tmin is shown in Fig 8. The obtained curve reveals three successive exothermic peaks. The DSC curve shows that the decomposition starts at 436*C. The first peak at S70*C and spans the range of 436-590'C. the second one at 603'C and spans the temperature range 592-610*C. and the third one at 630*C and spans the temperature range of 612«fl45*C. Comparing the obtained data with the reported ones of binary AJ-Mn alloys shows that there is a remarkable differences between the thermal stability of the binary a0oys{5-7] and the present one. The differential thermal analysis of the binary alloy Ai-15wt%Mn showing a single peak at 390*C and spans the temperature range of 32S4S0*C at a heating rate of 10*C/mta This peak is assigned to the crystallization of the i-phase into the equilibrium AlgMn. An additional peak is reported to appear with increasing Mn contents. For binary alloys contain Mn>B.6artW,17.2wttt) the DSC data show two overlapped peaks at 415 and 430*C(6> which separate into two sharp peaks at 466*20*C and 605*12*C at Mn contents >18at% (36wt%H5t These peaks are assigned to the crystallization of both tha i-phase and the t-ohase mto the equilibrium AlsMn respectivery[51. The transformation route of the metastable constituents of the as-matt quenched structure of the present alloy is observed to be completely different from that reported for the binary AM5wt*M» aUoyfJ] that shows a direct crystallization of the i-phss« into the ec^jilibrium AtyMn(5>7J. But the thermal analysis data supported by the X-ray results of the present alloy illustrate that the i-phase transformed firstly into two transition phases: the metastable AigMn and the t-phase. This transformation corresponds to the first peak at 570°C. Next the metastable AlgMn transformed into the equilibrium AtgMn at 603*C. which is the second peak. Finally the t- phase is shown to crystalline into the equilibrium AlgMn at 63S"C, which is the third peak. A lower crystallization temperature of Ihe t-pnase of 605±12* C is given for binary alloys[5I. In the light of the previously discussed results it can be deduced that the presence of titanium even at such a low content as 2wt%, thermally stabilizes the i-phase to decompose at higher temperature than that observed in the binary alloys. In general the DSC curve and the XRD results show the complexity of the transformations that take place dunng the annealing course of the as-matt quenched structure of the AM5wMMn-2wntTi alloy CONCLUSIONS: Formation and thermal stability of the i-phase formed in the as-melt quenched AI-15wtMn%- 2wt9tTi alloy » investigated Based on both the obtained and published results it can be concluded that 1- The i-phase is observed to form in the as melt-quenched rmcrostructure embedded in a matrix of a super saturated solid solution of the u-AI that is estimated to contain about 5 wt%Mr> . 2- Titanium is found to thermally stabilize the i-phase to start decomposition at 436'C Furthermore, the decomposition route is observed to be utterly different from that reported for binary alloys The i-phase decomposes into two transition phases, metastable AI6Mn -M2- Table I X-Ray Diffraction Data for As-Mc.lt Spun Ribbons Al-15Mn -2Ti Al-lfMn Al-(18-25)Mn Al-28Mri d,nm I/Io i-phase d,nn> 1/iu d,nm l/Io d,nm l/lo 0.3877 7 (110001) 0.2314 73 0.3830 7 0.3850 22 0.3343 5 (111010) 0.2175 22 0.3330 *> 0.3350 8 A. 0.2176 25 (100000) 0.2068 24 0.2168 100 0.2186 100 0.20C8 25 (110000) 0.2005 15 0.2062 75 0.2066 78 0.1492 3 (111000) 0.1497 3 0.2056 75 0.1496 11 0.1274 7 (101000) 0.1277 6 0.1479 7 0.1275 20 0.1101 3 (110010) 0.1206 5 0.1276 21 0.1102 5 0.1085 3 (200000) 0.1085 3 0.1106 6 0.1086 7 0.2324 100 0.2325 100 0.1086 5 0.2012 25 0.2020 20 0.1419 14 0.1427 20 0.1215 12 0.1215 5 0.1157 6 0.11C4 6 0.0923 6 0.0924 9 Note: All tlie italic reflections are due to Aluminum pivasc. The rest are due to the i-phase. -343- Table 11 X-Ray Data Reflection for Annealed Ribbons at 400°C for Different Intervales of Time 30 min. one hour tuo hours four hours d,nm I/lo d.ran I/lo d,nm 1/1° d,rm I/lo 0.6258 6 0.6294 4 0.2376 8 0.2374 19 0.414S 5 0.4149 4 0.2313 2 0.2246 5 0.2475 3 0.2485 2 0.2281 3 0.2251 S 0.2458 3 0.2377 6 0.2255 2 0.2239 9 0.2376 3 0.2297 1 0.2241 5 0.2218 5 0.2303 2 0.2250 4 0.2189 21 0.2186 19 0.2243 4 0.2191 19 0.2161 11 0.2167 18 0.2190 21 0.2162 11 0.2134 3 0.21S5 13 0.2165 17 0.2142 4 0.2083 17 0.2131 S 0.2090 B 0.2107 3 0.2067 24 0.2090 7 0.2070 30 0.2089 13 0.2048 9 0.2069 28 0.2055 1C 0.2072 22 0.2008 1 0.2053 16 0.2010 4 0.2056 10 0.2342 100 0.2007 20 0.1997 2 0.2011 S 0.2O2C 33 0.1904 4 0.192G 2 0.1921 1 0.1876 4 0.1910 2 0.2343 100 0.2339 100 0.2335 100 0.2027 23 0.2025 22 0.2022 28 Table 111 X-Ray Diffraction Data for Annealed Ribbons at Various Temperatures up to 620 C 3S0°C/1 hr 4S0°C/1 hr SSOt/SO min SS0°C/24 hrs 620°C/1 hr d.nml/l! o dvnm l/Io d,nm I/lo d,nm I/I. d.nri I/I 0.6188 5 0.6267 9 0.2483 6 0.6267 9 0.6267 8 0.5627 6 0.S336 4 0.2435 3 0.4357 11 0.4931 11 0.51E0 6 0.4141 5 0.2382 30 0.4454 7 0.4455 7 0.4130 7 0.2251 9 0.4304 3 0.4304 5 0.4309 4 0.3878 7 0.2372 14 0.2192 26 0.'41S3 4 0.4288 2 0.3564 4 0.2284 4 0.2170 30 0.3789 3 0.4141 2 0.2637 3 0.2239 6 0.2141 11 0.3288 7 0.3821 7 0.2642 2 0.2164 25 0.2075 41 0.2881 2 0.3523 3 0.2429 3 0.21C4 23 0.20SS 23 0.2627 5 0.3288 7 0.2374 5 0.20C6 35 0.2012 29 0.2539 7 0.303C 9 0.2323 5 0.2006 13 0.1881 4 0.2472 7 0.2623 8 0.2299 6 0.1878 3 0.2341 100 0.2428 3 0.2534 4 0.2166 27 0.2341 100 0.202G 26 0.2372 10 0.2338 12 0.2155 13 0.2027 24 0.2273 1G 0.2270 17 0.2148 C 0.224S 8 0.2195 9 0.2077 23 0.2219 11 0.2188 14 0.1968 3 0.2188 30 0.2155 33 0.2334 100 0.2159 28 0.2130 30 0.2022 34 0.2135 7 0.2081 9 0.2073 55 0.2076 56 0.1910 5 0.2006 6 0.1691 7 0.1910 9 0.1834 3 0.1691 3 0.1646 2 0.1646 3 0.2339 100 0.2331 100 0.2025 55 0.2025 41 -345- .^^V^.v^**d*. L W 11.111 ,_Ak/uJLi.w^._-Jik?'',1Ui.*..._.J,jj .vi.j.*wUL /vWuJ,..Jvfi u^ A*k&HmULualL ixuLuJ^iLL 10 20 40 60 80 86 28 (deq.) Fig 1 Changes in X-ray reflection profiles of as-melt spun Al-15Mn-2Ti alloy after annealing at 400°C for several intervals of time. a) as-melt spun, b) 0.5 hr, c) 1.0 hr, d) 2 hrs, and e) 4 hrs. -346- l.«enr o .S f \ ^KtMiiv .ffclu^. MJALL "'W^^L L JL/^ JJIjyjJll! 10 20 40 ^60 80 66 29 (deg.) Fig 2 Changes in X-ray reflection of as-melt spun Al-15Mn-2Ti alloy after annealing at several temperatures. a) at 350 °C/ 1 hr, b) at 400 °C/ 1 hr, c) at 450 •<:/ 1 hr, and d) 550°C/ 24 hrs. -347- • 0. 910 . _ »^—i •• 0. 936 t ° e c / ••«. 0, 92< —-*"*• \/~~ ,-* 500 *C TT 0.92 4 • 400 *C 0. 922 • 1 « 1 2 3 annealing tine / hr b 0.930 E /*"* S 0.926 / 0.922 . 1 _ A _.l l_.l_ _» . 0 200 400 600 600 «nncalln9 tcr»pcr«tur» / *C Fig 3 Changes in (331) aluminum interplanar spacing with: a) annealing time at 400 and 500°C, b) annealing temperature, holding time one hour. Fig 4 Microstructure of as-melt spun material -348- J50 H Fig 5 (a) Changes in Vickers hardness 250 of as-melt spun 200 400 600 annulling t«ap«r*tui« *C alloy with annealing temperature. b) Changes in Vickers hardness 350 of as-melt spun alloy with annealing 100 time at 550°C. 2J0 200 0 (0 120 100 240 *nn«allng tint / Bin. t 4J id 0) o S •o c 0) 35_u0- 400 450 500 550 600 650 temperature/°C Fig 6 DSC curve obtained at a heating rate of 10°c/min for as-melt spun material. -349- cind ifc*. ;vl".uso that ;irut!ly u anile;.r.od intj ,nc tqwlibiiuni Al6Mn at temperatures of 603 and ojb "C respectively Several stable intermc-tahcs (AI6Mn. (JlnsTi. and AI3Ti) arc detected in the ribbons after REFERKNCKS: 1- D.Shechtman, I. Blech. and J.W. Than, Phys Rev Lett.. 53(1984) 1951-53. ?- PA Bancae et al. Phys Rev Lett . 54(1985) 2422-25. ?- R.J. Schaefer et al. Metall Trans, 17(1986) 2117-25 4- D. H. Kim and B. Canter, Scripts Metall., 23( 1389) 1859-04. 5- K.F. Kelton and J.C. Holzer. Mat Sci. Eng., 99(1988) 489-92 6- K. Yu-Zhang et al. ibid. 385-88 7- F H Samuel et al. Metall. Trans., 17(1986) 1671-83. 8- D. Shechtman, R.J. Schaefer. and F.S. Biananiello, Metal!. Trans., 15(1984) 1987-97 9- R.D. Field and H.L. Fraser. Mat. Sci. Eng.. 68(1984) L17-L21. 10- L. Mondolfo. "Aluminum Alloys Structure and Properties", Butterworth, London, (1976). 11 - T. Ohshi, L Dai, and N. Funkstus, Metal! Ti iin:,, 17(19eG) 739-807. 12- D Shecntman and IA Blech, Metall Tians. 16(1985) 1005-12. 13- P. Villars and L.D Calvert ."Pearson's Handbook of Crystallographic Data for Intermetallic Phases", Vol. 2, ASM(1985), 1021. 14- M.A.T«ylor, Acta Metall.. 8(1960) 258-62. 15- K. Liitfe and W. Hume-Rotriery. J. Inst. Metals, 74(1947)521-24 16- E. Nes. S.E. Naess, and R. Hoier, Z Metall., 63(1972) 248. . 17- JCPDS Card File No. 37-1449 18- JCPDS Cord File No. 31-S4C 19- JCPDS Card File No 4-787 -350- Jb GrS°'/S2 Plastic Deformation Instability in Tension in Alloy 800 H M. E. Abd El-Azim and F. H. Haramad Metallurgy Dept., Atomic Energy Authority Cairo, Egypt Abstract Alloy 800H was tensile tested at 22. 650,750 and 800°C up to different strains al each temperature. The variations of the reduction of the cross-sectional area and of the strain along the gauge length of each tensile specimen were determined. It was found that al 22 and 650°C the jdastic deformation instability i.e. necking occurred at a strain at which maximum load was reached in agreement with Considcrc criterion Al 75(1 and K()0°C, (he material underwent steady slate deformation and necking occurred at a strain much larger than the strain at which maximum load was reached. Keywords Alloy 800H; plastic instability; necking; tensile test; reduction of area; strain; deformation temperature; steady state deformation; microstructure. 1. Introduction It is well known that in a tensile test nonuniform plastic flow occurs when the maximum load is reached . at which point the specimen becomes mechanically unstable and the plastic deformation becomes localized in a neck. Considere"(1885) gave the first quantitative treatment which indentificd the onset of necking with the attainment of a load maximum. At the maximum load the geometrical softening of the specimen due to reduction of the cross-sectional area becomes equal to the rate of work hardening. Hart (1967) and Campbell (1967) formulated detailed theories for the development of instability in the tension test in work hardening and rate sensitive materials. Halt's analysis was based on the growth of an imperfection in the cross- sectional area of the tensile bar. He postulated a criterion for a stable plastic deformation as follows. Suppose that at some stage of the test, there is a difference between two cross sections of the specimen. If this area difference becomes smaller or remains constant as the test progresses the deformation is stable. Otherwise, «he deformation is unstable. Campbell's treatment was based on the development of strain gradient in the (ensile specimen. Jonas ct al. (1976) introduced an additional instability criterion based on the rate of growth of a nonuniformity in strain rather than the cross-sectional area considered by Hart. Jonas ct al.(!976) analyzed the plastic stability of metals undergoing tensile loading in some details. Two types of external defects were considered : mechanical -351- damage involving work hardening and geometric faults due to machining variations . It was shown that, for variations in initial cross-section of similar magnitude, geometric faults led to much more rapid neck formation than mechanical defects. Even with geometric defects, however, the presence of faults of conventional magnitude did not lead to failure rapidly enough to form the basis for normal necking behaviour. Although Hart's and Jonas's analyses of the development of plastic instability in tension test are similar except for the choice of the relevant variable, the instability criteria are different. In terms of the dimensionless work hardening coefficient y, which is expressed by : y=l/<7(5o78e) (I) where a is the flow stress and e is the strain, and the strain rate sensitivity m, which can be expressed by: m = e/CT(8o-/8s ) (2) where e is the strain rate, the Hart's criterion for stable deformation is given by: yai-m (3) while the Jonas et al.'s criterion is given by: y > 1 (4) In a more recent paper Jonas and Baudelet (1977) claimed that the two criteria are not in conflict, but refer to different stages of neck development, depending on whether 8E or 8 A is taken as the relevant variable where A is the cross-sectional area of the tensile specimen. The aim of this research was to study the effect of deformation temperature (22-800 C) on the development of plastic instability in tension in Alloy 800 H. Alloy 800H (Fe32Ni20Cr) is a high temperature Fe-Ni-Cr austenitic stainless steel and is used a candidate material for the heat exchanger components in high temperature gas cooled reactors. 2- Experimental procedure Details of the test material are summarized in Table 1. Tensile tests were performed using Zwick mechanical tensile testing machine equipped with a three zone furnace. Specimen temperatures were measured by three thermocouples and maintained within ± 3 C of the nominal test temperature. Cylindrical tensile specimens of a uniform gauge length 34.75 mm and of a diameter 6.4 mm were used. A length of 30 mm of the uniform gauge length of each tensile specimen was divided by 7 marks into 6 equal intervals in order to measure the cross - sectional -352- area of the specimen at these marks and the strain which occurs in each interval after the tensile test. All strains reported in this research are true plastic strains. 3- Results 3.1. Reduction of area and change of strain In order to study the plastic instability of flow in the tensile test in Alloy 800 H, the reduction of the cross - sectional area at 7 marks , which were marked along the uniform gauge length (30 mm) of the tensile specimen before the tensile test and separated by equal intervals of 5mm length, was measured. In addition, the variation of strain along the uniform gauge length of the tensile specimen was determined by determining the strain which occurred in each interval of the 6 intervals between the 7 marks marked along the uniform gauge length of the specimen. Figures 1-4 show the variation of reduction of the cross - sectional area and of the strain along the uniform gauge length of the tensile specimens which were strained up to different strains at each test tcmpcralurc in the temperature range 22 - 800 C and at a strain rate of 1.2 x 10'3 s"1. These figures show also the maximum value of the reduction of the cross-sectional area which occurred in some specimens at a point between the 7 marks along the gauge length of the specimen. At room temperature (Fig. I), the specimens were strained up to 36.8, 37.7. 39.3 and 42.8 % strain, respectively. The average strain, at which maximum load was reached, was 34.7% as shown in Tabic 2. The specimens strained up to 37.7. 39.3 and 42.8 % strain show maximum in the reduction of the cross-sectional area along the uniform gauge length of the specimens (Fig. 1 (a)) accompanied with maximum in the strain (Fig. I (b) ) in the interval 4 for the specimen strained up to 37.7% strain and in the interval 3 for the specimens strained up to 39.3% and 42.8 % strain. The values of these maxima increased with increasing the strain up to which the specimen was strained. On the other hand the specimen strained up lo 36.8% strain does not show maximum cither in the reduction of the cross-sectional area or in the strain along the gauge length of the specimen. But it shows only a small difference (£ 6.9 %) between the reductions of the cross-sectional area along the uniform gauge length of the specimen. Figure 2 shows that at 650°C the tensile specimens strained up to 38.0.38.5 and 40.5% strain exhibit maximum in the reduction ofthecross-seclional area associated with maximum in the strain along the uniform gauge length of the specimens. The tensile specimen strained up to 35.7% strain which is nearly equal to the strain (35.8%). al which maximum load was reached, docs not show maximum cither in the reduction of cross-sectional area or in the strain along the uniform gauge length of the specimcn.Thcrc is only a small difference (2 3.7%) between the reductions of the cross-sectional area along the gauge length of the specimen. -353- Tabic ('.Conditions of test material. Material Alloy Form Ileal treatment Grain sue Klomciil analyM*. »«•/• manufact (pill) urcr «. Fe A7 o Ti Al S Allov VDM Co bars of Soluli.n 258 0.065 bal 30.3 20.05 0J2 0.30 0.003 800H Germany diamcicr treatmem at 25mm 1150° C for 1 h 1 : 1 Tabic 2:i£fi'ect of'(est temperature on average strain at which maximum load was reached lor Alloy 800M. Test temperature .Strain at maximum load Average Standard deviation 22 34.7 ±0.35 650 35.8 xO.74 750 22.3 ± 0.89 800 9.6 ±0.84 100 17 0 Alloy BOOH Alloy 800H £ •/, » uj t 100 A • 39 3 <. 3» 7 SO / \ IA o J6 e • co - c 70 - 22°C f:i2.io'3s-' 0 . ...i— _.i J— iiil IC 70 30 Distance olong the gcuge length.•*<* |r.i*fyo!S alon; ti-e gouge lengln ;ol to) Fig 1 Vorialion of the reduction ol the cross • sectioncl area dnd of trie strain oiong the uni'orm gauge l*ngth of the te.'.sile specimens of Alloy 300 -• sfsmed to fliflerent strains ot 22°C,la] Reduction z> crea . Ibl strain 90 Alloy 600 •< 80 Alloy 800 H t V. 70 * *0S 70 £ 38 5 ;> so 60 8 o 38.0 c SO ' f\ < 35.7 SO "o - / I 40 30 6 30 20 650°C 650° 10 • 6: 12.10'3 s"1 €:12.10"V 10 1 0 10 70 30 (0 SO ' 2 3 * 5 i Oislonce along the gouge length.mm tnlrvals along 'he gouge length la) ( Bl Fig 2 Variolion of the reduction of the cross - sectional area and of the strain along the uniform gauge length of the tensile specimens of Alloy BOOH strained to different strains ot 650°C.lol Reduction ot area . |bl Strain Allny ft'ifi M 90 • '.*< f, 80 • 34 ft 70 >- 71 (. * . 10 * —"" r M> • )0 ...-o o- -^ 20 v o~-<- 7S0°C 10 £;!.? -lO"3 S"' 0 t • i i i i CistCf'Ce olong i he gauge length,mm Intrvals otcnj the gauge length la I !b) :'9 } -'crraito'' ot the reduction ot the cross -sect'rnol area ar-c of *h? stroin along the uniform 3-uj* length o< the tensile specimens ol Ahoy 500H slrcmed 'o diflerprit slrnms at ~:0CC . 'aifieCuction ot oreo . !bl Strom l'O Alloy SO?"' Alloy 800H 800°C 120 £:i2iiO-3>-' 100 £ 60 5 60 40 20 _i_ J_ I 2 3 4 5 6 Distance along the gauge length.mm Intervals along the gauge length la) lb) Fig i Variation ot the reduction ol the cross -sectional area and ol the strain along the uniform gauge length ol the tensile specimens ol Alloy 800 H stroined to different strains ot 800°C . la I Reduction of area . Ibl Strain -356- At 750°C. the tensile specimen strained up to 35.8% strain, which is larger than the strain (22.3%) at which maximum load was reached, docs not show maximum either in the reduction or the cross-sectional area (Fig. 3 (a)) or in the strain (Fig. 3 (b) ) along the gauge length of the specimen. But it shows a small difference (£ 8.3%) between the reductions of the cross-sectional area along the gauge length of the specimen. The specimens strained up to a strain higher than 35.8% show maximum cither in the reduction of the cross-scctionni area or in the strain along the gauge length of the specimen. In additon , at 800°C, the specimen strained up to 43.6% strain, which is much larger than the strain (9.6%) at which maximum load was reached, docs not show maximum either in the reduction of the cross-sectional area (Fig.4(a)) or in the strain (Fig.4(b)), along the gauge length of the specimen. The specimens strained up to a strain higher than 43.6% exhibit maximum cither in the reduction of the,cross-scclional area or in the strain along the gauge length of the specimen. 3.2 Mctallographic examination Figure 5 shows the microstnicturc of the longitudinal section of the tensile specimen which was tested at room temperature, strained up to 39.3% strain and showed maximum in the reduction of the cross-sectional area as well as in the strain along the gauge length of the specimen. Voids can be observed in the microstructure around the fractured primary carbides, which were fractured by deforming the specimen up to this strain. Figure 6 shows the microstructtircs of three longitudinal sections of three tensile specimens strained up to 50.8, 43.4 and 34.5% strain, respectively, at 800° C. The microstnicturc of the longitudinal section of the specimen, which was strained up to 50.8% strain and showed maximum in the reduction of the cross- sectional area as well as in the strain along its uniform gauge length, shows (he formation of voids around the intra - and intcrgranular primary carbide particles (Fig . 6 (a)). On the other hand, the microstructurcs (Fig . 6 (b.c)) of the longitudinal sections of the specimens which were strained up lo 43.6 and 34.5% strain and did not show maximum cither in the cross sectional area or in the strain along the uniform gauge length of the specimen, show only the formation of smaller voids around the primary carbide particles. 4- Discussion Plastic instability of flow stress (necking) of metals undergoing tensile loading appears in the form of a difference in area between two cross - sections of the tensile specimen according lo Hart's (1967) concept of plastic stability or in the form of a variation in logarithmic strain or engineering strain according to Campbell (1967) and Jones ct al.'s (1976) concept. So the results shown in Figs. I - 4 were given in the form of variation of the cross sectional area and variation of the true (logarithmic) strain along the uniform gauge length of the tensile specimen. It appears that the two criteria of plastic instability are not clearly different. When there is a maximum in the reduction of the cross-sectional area along the gauge -357- .•••*v 4 V, • • ••'» /W ./' '•• /• ' • '1 KJ, • II I Ml I |l Tl ll.lll «[«» Fig.* Ojrtici! micropaph nl tlie longitudinal section of tensile specimen of Alloy 80011 strained at room temperature lo 39.3% strain (X200) Y7 7 J' •'. / V-rv • ) ,4 • ; t$Kt '7. * ;. »' ^ V'''&£~:U\ 6(b) 6{c) Fig. 6 Optical micrographs of the longitudinal sections of (ensile specimens of Alloy «00H strained at 80O"C • to dilTcrent strains: (a) 50.8% strain (XIOO) (b) 43.4% strain (XIOO) and (c) 34.5% strain (XIOO) -358- length of the specimen, there is a corresponding maximum in strain in the interval which shows maximum in the reduction or the cross-sectional area as shown in Figs. 1-4. However Jones and Gillis( 1978) pointed out that all of lltcsc criteria of plastic instability can be classified as rules for determining only the onset of instability, thai is, small instability. According to Considered (1X85) criterion, necking in uniaxial (ensile tcsl occurs at the strain at which maximum load is reached . This condition is snlisficd for Alloy 800H at room temperature (22°C)aiid at 65<>°C where necking occurs at a strain approximately equal to (he strain at which maximum load is reached (sec Table 2 and Figs. 1 and 2). At 22 and 650°C, the deformation behaviour of Alloy 800 H is less sensitive to the change of strain rale (Abd El-Azim el nl.. 1990) and (he material has well denned work hardening nearly over the whole strain. This work hardening appears in the form of three stages I, II and II! of hardening in the true slrcss-true strain curves at 22 and 65()°C as shown in Fig. 7. At 750 and 800°C, Considered criterion can not be satisfied . Necking occurs at a strain much larger than the strain at which maximum load is reached (sec Table 2 and Figs. 3 and 4). At these temperatures the deformation of Alloy 80UH is very sensitive to the change of strain rate and lest temperature (Abd El- Azim ct al., 1990; Abd El -Azim et al., 1992). The material shows three stages I, II and III of hardening up to 23% strain at 750°C and up to 10% strain at 800°C followed by steady stale deformation stage (IV) which extends up to 37% strain al 750°C and up lo 52% strain at 800°C at strain rate of 1.2 x 10 "3 s*1 as shown in Fig. 7. The steady state deformation occurring in hot tensile tests of Alloy 800H is temperature and strain rale dependent (Abd El - Azim ct al., 1990;Abd El-Azim ct al., 1992) similar to steady stale creep. The steady state stress i.e. the saturation stress in the hot tensile tests of Alloy 800H is temperature and strain rate dependent and the work hardening rale in this stage is zero (Abd El - Azim ct nl., 1990; Abd El-Azim ct al., 1992). The steady state stage in hot tensile tests results from the balance between work hardening and dynamic recovery processes (Mccking and Grinbcrg, 1979). The major portion of the steady state deformation in the hot tensile test of Alloy 800H at 800°C, which extends from i.bout 10 to about 43.6% strain, is a uniform deformation where deformation instability (necking) occurs at a strain larger than 43.6% strain. Estrin and Mccking (1980) found that deformation instability can be ruled out the steady state creep when the applied stress is smaller than 6 /N where 9 is the work hardening rate in stage II of hardening in the stress - strain curve and N is the exponent in power law expression for the steady stale creep. By applying this case to the hot tensile test of Alloy 800 Hat 800°C and at strain rate of 1.2 x 10"3 s"1. it was found that the value of the saturation stress in the steady state stage.in the hot tensile test, which is similar to the applied stress in creep test, is about 230 MPa and is smaller than the value of 6/N which is about 256 MPa. -359- Jonas and Baudclct's (1977) analysis of plastic stability, which was based on (lie criteria of Jonas ct al. (1976) and Hart (1967), was extended to include the generation by deformation of spherical cavities on the one hand and of planar cracks on the other. It was found that the generation of spherical cavities and voids or planar cracks during straining increases the tendency towards tensile instability by reducing the effective values of the work hardening coefficient and the strain rate sensitivity (Jonas and Baudclct, 1977). When the crack or cavity generation rate varies along the gauge length of the specimen, initial defects arc not required to initiate flow localization. When mechanical and machining defects arc present in addition to defect generation gradients, the latter will accelerate the rale of flow localization and thus decrease the fracture strain (Jonas and Baudclct, 1977). In the presence of several incipient necks, defect generation rate gradients can be responsible for eventual flow localization at a single site (Jonas and Baudclct. 1977). So the formation of voids due to the fracture of the primary carbides themselves by deformation as shown in Fig. 5 is responsible for the occurrence of necking at 22% in Alloy 800H. At 800TJ the formation of voids due to (he decohesion of the intra-or intcrgranular carbide particle-matrix interfaces (see Fig.6) is rcsponsablc for the occurrence of necking in Alloy 800H. Since Alloy 800 H underwent steady state deformation at 800*fc (sec Fig.7), the plastic instability (necking) is controlled by the steady state deformation. The steady stale lowers the stress concentration around the carbide particles because of the occurrence of dynamic recovery. These low stresses arc not sufficient to fracture the carbide particles but to cause the decohesion of the carbide particle - matrix interface as shown in Fig. 6. Also the steady state deformation delays the growth of voids to larger strains as shown in Fig. 6. During the uniform steady state deformation, there is no significant difference between the size of voids formed at the carbide particle - matrix interface in the specimen strained up to 34.5% strain in comparsion with that of voids formed in the specimen strained up to 43.6% strain. On the other hand, the voids formed at the carbide matrix-interfaces in the specimen which was strained up to 50.3% strain and showed deformation instability are larger than those formed in the specimens which where strained up to 34.5% and 43.6% strains and did not show deformation instability. It was found that steady state deformation firstly delays the nuclcation of voids at (lie particle - matrix interface to larger strains. Secondly, the strain rate dependence of the flow stress can stabilize flow and thus dalay the growth and coalescences of voids to larger strains (Ashbyetal., 1979, Asliby, 1981 and Gandhi 1983). Thus the steady state deformation delays the occurrence of necking in Alloy 800H to larger strains as found at 80(J°C, front the strain of 9.6% at which maximum load is reached to a strain higher than 43.6%. 5. Conclusions I- The plastic deformation instability (necking) occurred in Alloy 800 H, tensile tested at 22 and 650°C. at a strain nearly equal to the strain at which maximum load was reached. This condition is in an agreement with Considered criterion. -360- OL z \> 60 Fig.7; True-stress -true stroin curves of Alloy 800H at various temperatures -361- 2- At 750 and 800% Alloy 800H underwent steady stale deformation, which delayed the occurrence of necking to larger strains. At these temperatures necking occurred at a strain much larger than the strain at which maximum load was reached. 3- Plastic instability in Alloy 800H results from the formation of voids cither due to the fracture of the primary carbides during straining at room temperature or due to decohesion of the carbide particle - matrix interface during straining at 800°C. References Abd El-Azim, M.E. and Hanunad, F.H.. (1992). Mcch. Mater., Vol. 14, 33. Abd El-Azim, M.E.. Ennis, P.J., Schuster, H., Hammad, F.H. and Nickel, H., (1990), Jul - 2344 Report, KFA, Jiilich, Germany. Ashby,M.F.,(1981), Prog.Matcr. Sci.Chalmcrs Anniversary Vol., 1. Ashby.M.F.Gandhi.C. and Taplin, D.M.R., (1979), Acta Met., Vol. 27, 729. Campbell, J., (1967), J. Mcch. Phys. Solids, Vol.15, 359. Consider!*, A.. (1885). Ann. Pouts Chnussccs. No. 6. Vol. 9,574. Estrin, Y. and Mccking, H., (1986), In: Proc. of the 7 th Inter. Confer, on "Strength of Metals and Alloys", H.J. McQueen, J.P. Bailon, J.I. Dickson, J.J. Jonas and M.G. Akbcn, cds., Pcrgamon Press. Vol. 1.. 589. Gandhi, C. (1983) , In: Proc. of Confer, on " Flow and Fracture at Elevated Temperatures". R.Raj, ed., ASM, Metals Park.Ohio.USA, 83. Hart, E., (1967) Acta Met., Vol. 15,351. Jonas, J.J. and Baudclct. B., (1977), Acta Met., Vol. 25.43. Jonas, J.J.. Holt. R.A. and colcman, C.E.. (1976). Acta Met.. Vol. 24,911 Jones, S.E. and Gillis, P.P., (1978), In "Analysis, Modelling and Experimentation", S.S. Hcckcr, A.K. Ghosh and H.L. Gegel, eds.. New york, TMS - AIME, 46. Mccking, H. and Grinberg A.. (1979), In: Proce. of the 5th Inter. Confer, on "Strength of Metals and Alloys", P.Haascn, V.Gerold and G.K.Kostors, eds.. Pcrgamon Press. Oxford. Vol.1.289. -362- First International Spring School A Symposium in Advances in Material Science (SAMS 94) Cairo, 15-20March 1994 EVALUATION OF ABRASIVE WEAR RESISTANCE OF CARBURIZED AND PLASMA SPRAYED 15CrNi6 STEEL II. Nasr and S.A. El-Ghazaly Engine Factory, AOI, Helwan, Egypt. Abstract Low alloyed carburizing steel such as 15 CrNi6 is widely used in case hardeuable crank shafts, gears and extruding sleeves. The rates of abrasive wear Onnr .h ) at constant pressure were measured for both plasma sprayed and carburized 15 CrNi6 steel under laboratory and industrial conditions. The effects of heat treatment cycles on the rate of abrasive wear for that materials were investigated. Moreover, topographical study of their abraded surface was achieved. 1. Introduction Many parts of engineering equipment are subjected to different kinds of wear in dry or wet conditions, for instance in acronctic p;ins, crank shafts and gears both severe drag forces and abrosivc-errosive wear took place in those parts. Moreover, in food industries, extruding sleeves made of carburized steel grade (15 CrNifi) are widely used. The rate of abrasive wear created during mixing and extruding the agregaies of food paste against the carburized steel surface determines to a high extent the overall cost of the process. Plasma repair techniques have been widely developed for iribological applications in many fields' '•*'. However, precautions about the limits of mechanical properties of the plasma treated parts must be taken into consideration for specific application™. Meanwhile, material engineers have to be aware with the weaknesses of plasma repaired parts relative to compact steel materii The microstructure of carburized steel after heat treatment affects to a great extent the abrasive wear resistance, while the homogeneity and compactness of plasma sprayed layer determine the success of die operation'6'. Accordingly: some factors affecting the abrasive wear resistance of both carburized steel 15 CrNi6 and its plasma sprayed parts were investigated. Those factors comprise the microsiructurcs of the steel and sprayed layer changing wiihthe heat treatment cycle chosen, their bulk hardness and carbon percent of the surface layer. Structure of the abrad ;d surfaces is also discussed. -363- EXPERIMENTAL Carburizing steel grade (DIN 15 CrNi6) in hot-rolled normalized conditions was used in manufacturing extruding sleeves with the following dimensions: Dimensions in mm Length Diam. Teeth Width Teeth Depth Stationary cylinder 350 130 10 3.5 Rotating screw 370 115 4 10 The extruders were carburized to about 2 mm depth using salt bath or charcoal in box carburizing provided that the composition of the steel is: Element C Mn Si Cr Ni SP wt% 0.12/0.16 0.4/0.6 0.15/0.40 1.4/1.7 0.03 0.03 The hear treatment cycles used in both salt bath and solid carburizing were chosen as indicated in Fig. 1. PREHEATING PREHEATING 650°C/1 h 600°C/2h CARBURIZING CARBURIZING 900°C/5 h 880°C/8 h HARDENING IN HARDENING IN OIL OIL 840°C/12 min 820°C/10 min TEMPERING TEMPERING 180°C/2 h 120°C/2 h AIR COOLING AIR COOLING (a) (b) Fig. 1: Schematic representation of heat treatment cycles used in: (a) Salt bath carburizing, (b) Charcoal solid carburizing of 15CrNi6 steel. -364- On the other hand, defined samples of the same steel in normalized state (8 mm diameter) were plasma sprayed using M1000 plasma spray unit to about 02-0.5 mm layer on substrate using carbon-molybdenum iron powder (1.1% C, 3.2& Mo. balance Fc). The sprayed parts were investigated in as-spraycd and heat treated conditions. A special heat treatment cycle was used and is very near to that for alloyed high carbon tool steels. Pin or ring laboratory tribometcr unit was used to measure the rate of abrasive wear of plasma sprayed as well as carburized 15CrNi6 steel. The ring was chosen to be silicon carbide disc of grade 60. Amray 18301 scanning electron microscope aided with EDAX PU9900 analysing unit was used to evaluate the microstructure and phase analysis of the samples. RESULTS AND DISCUSSIONS Structure of Carburized and Plasma Sprayed 15CrNi6 Carbon-molybdenum, iron base powder was sprayed on 15CrNi6 steel pins in normalized state. The structures of both normalized steel and plasma layer arc shown in Fig. 2. The normalized steel has a fine ferritic-peariiic-pearliiic structure, while plasma layer has continuous structure with carbides imbeded in the matrix. Qualitative and quantitative analyses of carbide phases using Amray scanning microscope and EDAX unit are shown in Fig. 3 and 4. It is shown that mixed iron and molybdenum carbides either with high concentration of iron or molybcdnum, arc present. On the other hand, the structure of carburized-hardened and tempered 15CrNi6, Fig. 5, reveals the homogeneity of carburized layer in case of solid than that in salt bath carburizing. The suucture of the case (1-2 mm) is tempered manensite whiie lower bainitc is present beyond. The carbon content of the case layer is about 0.65-0.8% depending on the depth of carburizing layer. Abrasive Wear Measurement Pin-on-ring tribometer was used to evaluate abrasive wear resistance of both plasma sprayed or carburized 15CrNi6 steel under different degrees of severity. Silicon carbidc-60 rings were used in laboratory wear measurements of both materials. Fig. 6, illustrates volume loss of plasma sprayed pin samples of pressure of 30N and ring rotation of 30 rpm as function in time. Depending on the hardness ratio of material to abrasive ring, wear occurs with different rales. It is evident that wear in heat treated (hardened) plasma layer (about 63 HRC) is minimized as with as-sprayed on steel pins. However, further studies must be concentrated to overcome separation observed between steel substrate and plasma layer. Fig. 7. -365- m ^i«fj^^»i? (a) X500 (b) X200 Fig. 2 : Microslruclures of both plasma sprayed (o) and normalized (b) l5C«Nii sleel. -**' T •- J»: Coibide (1) Carbide (2) Mo B6.9X 50.8X Fe 3.1X 40.2X C Dolonce Balance Fig. 3 : Carbides morphology ond composilion in plasma layer as revealed by SEM-EOAX unit. -366- li-tcp-io t 7: 10 ja e 3 A x READY RATE- 2 3 9 3 C P B TIME- taoLSCC PS- io ai •c N r . PAST- a O O (. S C C A •• a A i o PCAZMA- O O 1 B / 1 HJC« • • Point (1) • . J ._, -, . ... I,„ . peKB .i. !^ - • 2.00 I 6.00 S.OO a c N T g s K e v tO*v/en A COAX 1 Z • S a P - 9 0 t 7 • 2 2 : 3 4i E D A X a e A • V R A T E - 4 ; c P S T X K E - i U o I S E C F s.- s 9 a 9 c N T P n S T - > * o L S E C A - S A X o p u A Z M A - > o o > s / 2 FeKa Poinf (2) 1 f*iL« ' A . ... ^—~ ! <^» , •* -/ 2.OA.0 . , A.00 6.00 ,A B.00 IBaCNT 3.B2KEV :0*V,'cr> A COAX Fig. 4 : Distribution of Mo and Fe in carbides (I) and (2) revealed in plasma layer. -367- Soil Bath Coiburixed Steel, XIOO Solid Caiburizerf Sleet, XIOO Fig. 5 : Depth of carbuiized layer in !5CiNi4 steel. Force » 30 N Distance = 30 rpm SiC ring 40 i0 ?0 120 150 170 200 Ablation Time, (min.) Fig. & i Rale of abrasive wear measured for plasma sprayed lover In dilferenr conditions. -368- X600 XIOO Fig. 7 : Microttructurcs of plosmo lover and steel .sub*trat after lical Ireulcncnl. Normalized I.? Hardened 6 Fss I5CtWi4 EPS. Hardened 35CrNi4 1 Jo.8| Sail Bolh c Cacburized J^!v> )5CcNiA 0.6] 0.7BX C Solid Cutburized l5Cr-Nii 0.2 - 100 200 300 400 C.J -.1.0 1.5 2.0 Time, hrs. Time, hrs. Fig. 8 : tnd-jstrial and laboratory measurements of obrailve wear for ISCtHil in different conditions. -369- The industrial performance of carburized 15CrNi6 in food extrusion sleeves was measured in volume loss of the rotating screw teeth after many hundreds hours operation. Fig. 8, shows the relation between volume loss in rotating screw teeth and operation time as function in the carbon content of the carburized surface layer. On the other hand, laboratory measurements of abrasive wear observed on carbtirizcd slccl pin samples at different conditions were also achieved. Fig. 8. Abrasive Wear Appearance Thc.appcarancc of abraded surfaces of both plasma sprayed and carburized 15Cri6 steel gives valuable indications about the wear mechanism. In case of normalized and carburized steel (pins and sleeves), deep grooving is observed due to severe abrasion conditions, Fig. 9. While the topography of plasmaspraycd abraded surface is different, hence the softer phase is being abraded first and hard phases like molybdenum and iron carbides arc hollowed out from the surface. Fig. 9. That is important on selecting materials to suit specific applications. CONCLUSIONS The industrial and experimental investigation achieved in this article to reveal the relationship between the microstructurcs of carburized and plasma sprayed I5CrNi6 steel illustrates the following: 1. The abrasive wear rate of plasma layer (Fc-l.l%C-3.2%Mo) sprayed on 15CrNi6 steel has decreased drastically after heat treatment, however, care must be taken lo overcome interface cracking after heal treatment. 2. 15CrNi6 steel parts carburized using solid process have lower abrasive wear rate than that carburized in salt bath due to the homogeneity and compactness of the surface layer. 3. The appearance of the plasma layer abraded surface reveals the presence of hollow points due to the removal of nonhomogencous carbides of the matrix. ACKNOWLEDGEMENT The authors arc greatly indepted to Prof. Adcl Nofal, Head of Physical and Mechanical Section, CMRDI, Cairo, and Dr. Mohammed El-Scmary, Chairman of Engine Factory, A.O.I., Cairo for giving the many facilities and valuable discussions during this investigation. Sincere thanks arc due to BERZI Food Company in Cairo and the Physical Metallurgy Department of Miskolc University, Hungary for the unlimited support during application and evaluation of mair-ials used in this article. -370- Plasmo Sproyed Solid Cutbufiied Fig. 9 : Ab.uded surface, appearance of solid co.bu.ized and plasmo sprayed l5CfNi£ sled. -371- REFERENCES 1. Mayer, H.C. ci al., "Hcrstcllung von VcrschlciBschutzschichtcn durch DiffusionsschwciBcn und Elcktroncnstrahlumschmcl7.cn und dcrcn Eigcnschaftcn. Tribologic 9, Springer-Vcrlag, 77-156, (1985). 2. Kvcrncs, L. "Use or plasma spray technique in Diesel engine". Thin Solid Films, Vol. 53, No. 2 (1978). 3. Uctz, H., "Abrasion und Erosion", Carl Hanscr Vcrlag (1986). 4. Obcrlandcr, B., "Mctallographic studies of plasma sprayed duplex system". . Metallography 16, 117-115 (1983). 5. "General Inspection Requirements after Plasma Spraying", Socicta Nationale d'Etude ctdc Construction de Matcurs d'Aviation (SNECMA), F0301, DMP14- 010.284(1980). 6. Schulicr, P., "Mctallographische Utcrsuchungcn and Plasmagcspritzcn Doppclbcscichttungcn", Soundcrband No. 14,237-252 (1983). r^fcr- ibo/ibiLj First International Spring School & Symposium on Advances in Materials Science (S/\MS 94) 15-20 March 1994 CIIAKACTKKI/.ATION OK IIKTKKO-KIMTAX1AI, STIUKTUUKS BY X-KAYS AND IXKCTKON MICKOSCOI'Y NADIA A. EL-MASRY Materials Science and Engineering Department North Carolina State University Raleigh, N.C. 27695-7916, Fax (919) 515-3027 Abstract The performance of any thin film electronic device relics on the materials quality of the grown epitaxial layers. We will demonstrate the use of high resolution electron microscopy and x-rays in the evaluation of cp'iaxial thin film electronic materials. Hctcro- epitaxial Thin film* such as Cc02/Si, and InGaP/GaAs Mill be used as examples of materials characterized by these techniques. Introduction Advances and new applications in thin film electronic devices demand the heteroepitaxial growth of dissimilar materials. Hetero-epitaxy of single crystal - single phase thin Him faces many challenges. The following arc some cf the problems that faces hetcro-cpitaxy: to establish the growth conditions that produce an epitaxial single crystal and not polycrystalline films; overcome the lattice mismatch between the film and the substrate that lead to defect generation and poor device performance; minimize the thermal mismatch between the film and substrate that generates lattice defects during the cooling down from the growth temperature to room temperature that occurs in several systems such as GaAs/Si. Crystal mismatch is also a challenge where the grown film has a different crystal structure than the substrate crystal type such as hexagonal versus cubic or ionic crystal versus covalem crystal, e.g.," SiC/Si or CcOz/St, respectively. An intcrfacial reaction can occur if die film and substrate species tend to chemically rcaa or have the same chemical affinity with a common species or a diffusing component from or to the film, cg-CeOj/Si, II-VI compounds on GaAs. Also thin film growth of non miscible ternary alloys can face the problem of phase separation due to the themodynamical stability of the binary components that constitutes the ternary alloy, e.g., InGaP/GaAs substrates. From the above one can realize that the structure characterization plays an important role in identifying these problems and provides the evidence for its solution. In this work we will present two case examples to demonstrate the role of the structure characterization by means of x-ray diffraction and electron microscopy in identifying some of the above problems. Case 1: Cerium Oxide on Silicon Substrates The Silicon on insulator (SOI) is a premising structure for high speed devices '. The SOI structure requires thin epitaxial Si layer with low defect density ind a buried insulator with high endurance voltage for the expected performances 2. Various types of SOI structure were proposed and demonstrated. They include lateral growth from amorphous Sp, rcgrowth of amorphous Si by laser annealing4, separation by implanted oxygen (SIMOX)^ and direct wafer bonding0- \ The idea of using crystalline insulator seems to be preferable from the view point of crystalline quality and thickness control of the epitaxial Si layer on the insulator8. Recently, Ce02 has gained much attention as a pronising crystalline insulator due to its cubic CaF2 structure with a closer lattice parameter to Si than that of CaF2^' ^- The lattice mismatch estimated to be 0.35% '. The CeC>2, with a dielectric constant of 26, is also attractive as a capacitor material for the dynamic random access memory ". Potential application will rely on better understanding of CCO2/S1 interface and its quality control. The interfacial structure of Ce02/Si (111) is investigated in detail by high resolution electron microscopy and Auger electron spectroscopy. Based on observation and analysis, post annealing in oxygen atmosphere is carried out to modify the structure to Ce02/Si02/Si (111) with the expectation of improving the insulating properties of the film for SOI applications. The interface structure of the samples were investigated by high resolution electron microscopy (Topkon ATM 002B :HRTEM ) using cross sectional thin foils. The HRTEM was operated at 200 KV with the spherical aberration of 0.4 mm. X-ray diffraction method was used to identify existing phases of Cerium oxide on Si (111). Auger electron spectroscopy (JEOL JAMP30: AES ) was also employed to investigate in- depth profiles of elements in the samples. For the analysis, KLL transition was selected for Si and oxygen. For Ce, MNN transition was used to avoid signal overlapping in spectra. RHEED indicated, during the growth, the streaky 1 * 1 patcm during the Ce02 deposition. This means the existence of single crystal CaF2-type Ce02 on the Si (111) substrate. Figure 1 shows a HRTEM image of the intcrfacial structure of an as-deposited Ce02 film grown at 750 °C on a Si (111) surface. Between the CcC»2 overlayer and the Si (111) substrate there exist two distinct amorphous layers with dark and bright contrast respectively. The observed single crystal lattice image of the icp layer corresponds to the CaF2-type CeC»2. The CeC>2 has the B type orientation, which is a twin with 180° rotation, to the Si (111) substrate* K The dark contrast in the crystalline CeC<2 region is thought to be a strain field. The thicknesses of the dark and bright amorphous regions in Figure 1 are 6 nm and 4nm respectively. The total thickness of the amorphous regions is 10 nm, which is too thick to be considered a native oxide. Figure 2 shows the interface structure as modified by post annealing at 900 °C for 50 min. in O2 ambient This figure indicates that the dark amorphous region has disappeared and the thickness of the bright amorphous region increased. The CeC»2 layer still maintained the B-type orientation to the Si substrate. Figure 3 (a) shows a depth profile of Ce, Si, and oxygen in the as-deposited sample from AES analysis. In Figure 3 (a), from the surface to the interface, the cerium and oxygen profiles are constant At the interface region between Ce02 and Si substrate, the graded slope of the oxygen profile can be seen. This indicates that a concetran'on gradient of oxygen exists in the'interfacial region. In the post-annealed sample, as shown in Figure 3 (b), the cerium and oxygen profile are also constant from the surface to the interface. However, the slope of the oxygen profile at the interface is sharply reduced, and is followed by a short plateau, indicating the existence of another phase. From this profile, it is specualted that the Cc02 in the post annealed sample has an abrupt interface in stoichometry with the underlying SiC»2 phase. On the as-deposited sample, the C-V curve showed a break down voltage at 3.0V. The deep depletion was observed in the positive bias region due to short minority carrier lifetime, indicating a high density of recombination centers, or current leakage in the -374- insulator. However, on (he post annealed sample, the breakdown voltage was extremely improved, resulting in the achievement of the break down field above lMV/cm in this sample. Recently Tye reported an improved breakdovn field of the 9.0 MV/cm in the Ce02/Si02 bilayer structure12. In addition, a reduction of interfaci* states and a disappearance of the deep depletion was observed. As shown in Figure 1, single crystalline Ce02 was observed on the amorphous layers. The CeOi crystal seems to inherit a crystallographic information from the substrate because the B-type oriented CeOi single crystal is observed. From this, it is speculated that a reaction has occurred during tlit cpitaxi.il growth of CcC" on Si substrate. According to thermodynamic data'3, there is a possibility of a reaction expressed by: 4Cc02+Si=2Ce2Q3+Si02 AHf=-46.4 Kcal/molc. (1) where, AHf is the heat of formation for the reaction. During growth, oxygen atoms in theCe02 layer are released from their lattice sites, resulting in the formation of an oxygen deficient CeO* region. It is possible that this oxygen deficient CeOx layer corresponds to the dark amorphous layer in Figure I. Oxygen atoms from the Ce02 layer diffuse into the Si substrate, forming amorphous SiCVj. This S1O2 layer contrast, as observed byHRTEM in Figure 1. The proposed oxygen deficiency at the interface coincides with the graded oxygen profile established by AES analysis at the interface in the as-deposited sample. A driving force for this reaction may be reduction of the strain field at the Ce02/Si interface, amorphizing process due to strain field reportedly occurred in Si '4. Vacancies .which are generated in CeOx region due to the oxygen out-diffusion, may work effectively to reduce the activation energy for CeOx to be amorphized The oxidation at 900 °C in oxygen atmosphere altered the structure to that of Ce02/SiQ2/Si (111). The recrystallization of amorphous CeOx seems to occur from the Ce02 layer to the underlying CeOx layer by oxygen supplied during annealing. Also, the S1O2 grew thicker due to oxygen diffusing to the Si substrate. At this Ce02/Si02 interface, the s'oichometric change should be abrupt. The thicker Si02 region in the post-annealed sample is thought to be expressed as the short flat region in the oxygen profile shown in Figure 3 (b). The proposed reaction and recrystallizing process is schematically illustrated in Figure 4. In summary, the interface structure of Ce02/Si (111) was investigated by HRTEM and AES. At the Ce02/Si interface, a reaction occurred forming an oxygen deficient amorphous CeOx layer and a Si02 layer. Annealing in oxygen atmosphere caused the recrystallization of Ce02 and increased the thickness of the amqrphous Si02 layer. This structure was proved to have a low density of intcrfacial states and a high breakdown voltage. Case 2: Low Temperature Growth of InGaP/GaAs Low temperature (LT) growth of GaAs thin films by molecular beam epitaxy (MBE), at - 200°C, exhibit good crystalline quality and extremely high resistivity, and thus can provide excellent device isolation'5. The low temperature growth work was extended to other arsenic(As)-based compounds, such as InGaAs and AlGaAs, and to phosphorous(P)-based materials such as InP, GaP, and InGaP. LT InGaP can be -375- extremely useful since it can grow lattice matched to GaAs substrates. In the As-based compounds the low temperature growth causes excess As incorporation (-1 at%) in the lattice. The excess arsenic precipitates by annealing and causes the high resistivity effect in the film 16, The question in the P-based materials is wether excess phosphorous and phosphorous precipitates can exist and exhibit similar properties as the As-based materials. In this study we will discuss the structure characterization of LT InGaP, MBE grown, by Double-crystal x-ray diffraction (DCXRD), cross-sectional transmission electron microscopy (XTEM), and ultra-high resolution scanning transmission electron microscopy (STEM). DCXRD showed that the LT InGaP films had -47% In. For the sake of convenience, the ternary compound Ino.47Gao S>P is hereafter simply written as InGaP unless otherwise stated. Two LT structures were grown in this work for structural characterization of LT InGaP films. The first structure had only a single LT InGaP layer (200°C, 1.5 mm) grown on a GaAs buffer layer (550°C). Two Bragg peaks were observed for the DCXRD (400) reflection with the cpilayer peak located at a higher Bragg angle than the substrate peak, as shown in Fig. 5(a). This indicates that the LT InGaP film is not lattice matched to the GaAs substrate and its In composition is lower than the InGaP- GaAs lattice-matched composition of -49%. The narrow full width at half maximum (FvVHM) of the epilayer peak (20 arc sec) indicates the high quality of the crystalline epilayer. Annealing this sample at 600"C for 1 hr caused the angular separation between the epilayer and the substrate peaks to increase by 20 arc sec, as seen in Fig. 5(b). This observation indicates a slight reduction of the lattice parameter of the epilaycr after annealing. As to the lattice relaxation as a result of annealing, this is similar to the LT GaAs case. However, annealing LT GaAs at 600°C for -10 min results in a complete removal of the lattice mismatch while annealing LT InGaP at the same temperature for 1 hr results in only slight relaxation of the expanded lattice. This is manifested by a small peak shift of the annealed film to a higher Bragg angle. In order to compare the structural properties of the LT InGaP and the InGaP epilaycrs grown at higher temperatures (~480°C, hereafter referred to as HT InGaP), a second structure was grown, as shown in Fig. 6. During the growth of both HT and LT InGaP epilayers, a fixed PH3 flow rate of -4 seem was maintained It should be noted that there was an interruption of the PH3 flow into the growth chamber after the HT growth was ended and the substrate temperature was lowered to -300"C. The In compositions of the LT and HT InGaP epilayers were assumed to be identical by maintaining the same source temperatures during the layers' growth. DCXRD analyses from (400) reflection planes were performed on both the as- grown and annealed samples. TOT the as-grown samples, three (400) Bragg peaks were observed: the substrate, LT InGaP and HT InGaP peaks, respectively, as shown in Fig. 7(a). Upon annealing at 600°C for 1 hr the LT peak was found to be the only one shifted, as shown in Fig. 7(b). The observation that the annealing did not have any effect on the HT peak position was as expected. The angular separation between the LT and HT InGaP reflections was 70 arc sec before annealing and 50 arc sec after annealing. Therefore, the angular shift of the LT peak towards the HT peak after annealing was 20 arc sec, which was in good agreement with that of the first grown structure previously discussed. This relative shift, corresponding to the incomplete removal of the lattice-mismatch by 0.01%, might be caused by redistribution of the excess P atoms in the epilaycr lattice due to the annealing. The nature of the redistribution is not yet known at this time. Considering the HT InGaP as a "substrate", the LT InGaP/HT InGaP structure was found to have DCXRD characteristics similar to those of LT GaP/GaP'7. The fact that the LT peak was situated -37d- on the left side of the HT peak confirmed that the LT InGaP had a slightly larger lattice parameter than that of the HT InGaP {by 0.05% for the as-grown LT cpilayer). This lattice expansion can be explained by the presence of the excess P atoms incorporated into the epi layer due to the nonstoichiornetric growth. Both analytical STEM and DCXRD approaches revealed that the LT InGaP had -0.5 at.% more P atoms than the HT InGaP. The details of this determination wil 1 be published elsewhere. An XTEM micrograph (Fig. 3(a)) from the second structure indicates the presence of phase separation in LT InGaP epilayers. This is manifested in the form of a "precipitate- like" micTostructure. The precipitate-like phase separation (-800 A in diameter for the layer grown at -200°C) appears to decrease in volume as the grov/jn temperature increases. Annealing, however, seems to have a slight influence on the phase separated microstructurf- The phase separated areas were crystalline and strained as implied by the moire fringes shown in Fig. 8(b). Analytical STEM analyses indicated that the phase separated areas in the LT InGaP layer had a slightly higher In composition than the LT matrix. The HT InGaP showed a distinctive feature of slight decomposition. Both the phase separation in the LT InGaP layer and the slight decomposition in the HT InGaP layer might be explained by the spinodal decomposition usually observed in growth by liquid phase epitaxy (LPE) and metal-organic chemical vapor deposition (MOCVD). Spinodal decomposition causes the appearance of a two-phase microstructure with different In compositions at the growth temperature. The temperature-dependent resistivity and Hall effect measurements on the as- grown LT InGaP film revealed n-type conduction with a resistivity of 9x10s Wcrn and mobility of 120 cm2/V s at 296 K. It is necessary to point out that the LT InGaP remains semi-insulating before and after annealing. In conclusion, the LT InGaP showed a lattice expansion of 0.05%, as determined by DCXRD relative to the HT InGaP grown at 480°C with presumably the same In composition. The knice expansion can be explained by the excess Patoms of ~0.5 a\.7o present in the LT layer with respect to the HT layer. The lattice parameter of the LT InGaP did not return to that of the HT InGaP after annealing at 6O0°C for I hr. However, a slight reduction of the lattice expansion by 0.01% did occur after annealing, the possible explanation of which may be the redistribution of the excess P atoms. XTEM showed no formation of P-prccipitatcs after annealing, but did show phase separation in the as-grown LT InGaP epilaycr, manifested in the form of a "precipitate-like" microstructure. As with LT GaP17, annealing is not required for the LT InGaP films to be semi-insulating. This work was supported by the Office of Naval Research. References 1. For example, T.Nishimura, Y.Akasaka, and H.Nakata, in the Silicon-on-Jnsulaior : Its Technology and Application .edited by S.Furukawa, ( KTK Scientific Publisher, Tokyo, 1985) p263-268. 2. for example, J.C.Strum, Mat.Rcs.Soc.Symp.Proc.Vol 107 ,295 (1988). 3. Y.Kunii, M.Tabc, and k.Kajiyama, J.Appl.Phys., 54. 2847 (1983). 4. A.Ogura, and H.Tcrao J.Appl.Phys.,62, 4170 (1987). 5. K.Izumi, M.Doken, and H.Ariyoshi, Electron.Letts.,1 4, 593 (1978). 6. J.B.Lasky, Appl.Phys.Lett., 48. 78 (1986) 7. M.Shimbo, K.Furukawa, K.Fukuda, and K.Tanizawa, J.ADpl.Phvs., 60, 2987 (1986). -377- 8. H.hhiwara and T.Asano, Appl.Phys.Lett., 4 0, 66 (1982). 9. T.Inoue, Y. Yamamoto, S.Koyama, and S.Suzuki, Appl.Phys.Lctt., 5 6, 1332 (1990). 10. M.Yoshimoto, H.Nagata, T.Tsukahara, and H.Koinuma, Jpn.J.Appl.Phys., 2 9, LI 199 (1990) 11. R.T.Tung, J.M.Gibson, and J.M.Poate, Appl. Phys.Lett., 42, 888 (1983). 12. L.Tye, T.Chikyow, N.El-Masry and S.M.Bcdair (to be submitted to Appl. Phys. Lett.). 13. in the Metallurgical Thermochemistry edited by O.Kubaschewski and C.B.AIcock (Pcrgamon, New York, 1979) p278. 14. K.Minowa and K.Sumino, Phys.Rev.Lett.. 6 9, 320 (1992). 15. F. W. Smith, A. R. Calawa, C.-L. Chen, M. J. Manfra, and L. J. Mahoney, IEEE Electron Device Lett. 9, 77 (1988). 16. M. R. Melloch, K. Mahalingham, N. Otsuka, J. M. Woodall, and A. C. Warren, J. Cryst Growth 111, 39 (1991). 17. J. Ramdani. Y. He, M. Leonard, N. El-Masry, and S. M. Bedair, Appl. Miys. Lett. 61,1646(1992). -378- Figure 1 Lattice image of an as deposited sample from < 1 !0> direction. Almost single crystalline Ce02, amorphous CeOx (dark contrast .egion) followed by S1O2 (bright contrast region) were observed on the Si nil) substrate. Ce02 had a B- type orientation to the Si substrate. -379- Si02 Figure 2 Lattice image after thr rwst annealing in oxygen atmosphere from <110> direction. The dark amoipnous region disappeared and the bright amorphous region grew, indicating the regrowth of Ce02 and growth of S1O2. -380- AES analysis (as deposited sample) 70000 60000V- CcOx+SiO B 50000 40000 30000 20000^ 10000 Silicon 0 i • 'f t'iguro 3 lii-dcpdi profile of Cc, Si and oxygen in llic as deposited sample and die postannclacd sample analyzed by Auger electron spectroscopy, (a) In the a-, deposited sample, the gentle profile of oxygen at the interface indicated that a oxygen deficient GcO* region exists. AES analysis (after oxygen anneal) Figu re 3 CCOJ 1) initial epitaxial growth 2) reaction starts at the interface otveen 00 QO CCOJ SiO; Si(III) 3) Growth of amorphous CcO.« and SiO: •J) Regroivth of CcOi and growth ofSiOl Figure 4 Schematic illustration of the intcrfacial reaction. 1) At first, epitaxial Cc02 grows. 2) During the growth of CcQ2, the rcaciion starts at the interface. 3) After the deposition, the sample has a Ce02/CeOx (amorphous)/Si02(amorphous )/Si (111) structure. 4) After the annealing in oxygen ambient, the structure changes to Cc02/Si02/Si(lll). -383- GaAs substrate LT InGaP CO Z Z (a) >—i >< (b) JL -400 -200 0 200 400 ANGLE (arc sec) Figure 5 DCXRO from (400) reflection planes of the first structure described in the paper (a) as-grown and, (b) annealed at HT InGaP 4S0°C 1 urn GaAs buffer 550°C 1 urn SI LEC GaAs (100) substrate Figure 6 Schematic of the multilayer sample with the LT InGaP film grown on top of the HT InGaP film. -385- GaAs substrate LT InGaP > HT T.nGaP w z < i 200 200 400 ANGLE (arc sec) Figure 7 DCXRD from (400) reflectionplane s of the sair.nle shown in Fig. 2: (a) as- grown, the angular separation between the LT in'GaP and tlie HT InGaP reflections is 70 arc sec; (b) annealed at 600°°C tor 1 hr, the angular separation between the LT InGaP and the HT InGaP reflections is 50 arc sec. -386- AS •Sr3<* •HB, »»« V. \ Figure 8 XTEM micrographs obtained from die multilayer sample shown in Fig. 6 illustrate: (a) die layer structures and the phase separation, manifested in the form of a "precipitate-like'' microstructure, in the LTInGaP layer, and (b) the moire fringe in one of. the phase separated areas, as indicated by the arrow. -387- First International Spring School & Symposium on Advances in Materials Science (SAMS 94) 15-20 March 1994 An Overview on thin films E3ir & & si xr e c3. toy Chem i era. 1 Vapor Deposition I. D. Abdelr^zekf , Y. K. Af ifi1 , and S. Kr-avvroHrynfjRi2 1. Metallurgy Department, Nuclear research Center, A.E.A, Cairo, Egypt. 2. KPA, Juelich, Germany ABSTRACT Chemical Vapor Deposition, (CVD); involves the formation of a solid thin layer on a heated substrate surface by means of chemical reaction in gas or vapor phase. CVD techniques have expanded continuously and developed into the most important method for producing films for solid-state devices. CVD is considered to be the major technique for preparing most films used in the fabrication of semiconductor devices and integrated circuits. It has advantages such as the versatility, compatibility, quality, simplicity, reproducibility, and low cost. CVD has some disadvantages of; the use of comparatively high temperatures in many processes and chemical hazards caused by toxic, explosive, or corrosive gases. Chemical vapor deposition processes can be classified according to the type of their activation energy into thermally-activated CVD, plasma-enhanced CVD, laser-induced CVD, photochemical CVD, and electron-beam assisted CVD. In this paper an attempt is made to present all aspects of CVD equipment design and the variables affecting the deposition rate. Finally the preparation requirements and the application of CVD films are also summarized. l. nsrrFeorxjcrrioisr Coating is desirable, or may be necessary for -389- variety of applications; including economics, material conservation, unique properties, or the engineering and design flexibility which can be obtained by separating the surface properties from the bulk properties [1,2]. Film, coating, can be defined as a near-surface region whose properties are different from those of the bulk of the material. Because of the high technological importance of thin films, a wide variety or preparation techniques is presently available. There are different schemes for classification of coatings; a) Based on the processing techniques, coatings are classified as cladding, hot dip coatings, pack cementation coatings, electroplating and vapor deposition coatings, (chemical vapor deposition, physical vapor deposition) [2] . These techniques may also be classified according to the film formation environment; electrolysis (electroplating, electroless plating, and electrolytic anodization); vacuum (vacuum evaporation, ion beam deposition, molecular beam epitaxy, hot wall epitaxy, and ion implantation); plasma (sputter deposition and ion plating); liquid phase (liquid phase epitaxy); solid phase (solid pFase epitaxy); and chemical vapor (substrate chemical vapor conversion and chemical vapor deposition) [3,4]. b) Based on the extent of interdiffusion between coating and substrate, coatings may be divided into three categories, namely; diffusion coatings, overlay coatings, and ceramic barrier coatings [2]. Some of the above techniques will be discussed in this chapter: -390- 1.1. Diffusion Coating The preparation of a diffusion coating involves the deposition on and simultaneous diffusion into the substrate metal, of one or more elements from an external source to form a protective surface alloy rich in deposited elements. Diffusion coatings are carried out in separate stages [2]. a. Overlay coating An overlay coating is one in which an alloy is deposited on a substrate material with only a small amount of interaction or interdiffusion. Hot isostatic pressing or bra2ing, plasma spraying, and electron beam evaporation are techniques used to achieve overlay coating [2]. b. Ceramic barrier coating Ceramic coatings are normally complex silicates with metal or oxide additions [2]. c. Electrolytic deposition techniques Electroplating is a method of obtaining thin layers in an electrolytic solution under the action of the electric current, based on reducing the metallic ions at the cathode by introducing external electrons [2,4,51. d. Electroless plating Electroless plating consists of the continuous formation of a metallic layer as a result of chemical reduction occurring in a plating solution without applying an external current source [2,4,5]. 1.2. Vacuum Deposition Techniques a. Vacuum evaporation Vacuum evaporation consists of vaporizing the -391- source material by heating it resistively, inductively or with an electron gun in vacuum, followed by its vapor recondensation as a thin film. [3,4]. b. Ion beam deposition Coating by means of ion beams can be obtained by depositing the material directly from an ion beam containing the desired elements or compound, or indirectly by bombarding the target with external ions. This technique has the disadvantage of the need for a high vacuum with a low growth rate which have so far prevented widespread use of this method [3,5]. c. Ion implantation Ion implantation represents a method of obtaining thin layers which include atoms of the substrate. It is a process in which suitable ions are injected at a certain depth beneath the substrate surface by bombarding it with high-energy accelerated ions. A thin layer is formed on subsequent thermal annealing of the substrate which facilitates the chemical combination of implanted and substrate atoms. Ion implantation provides another method of forming thin films which incorporate surface atoms such as Si02, SijNj, and SiC [3]. 1.3. Plasma Deposition Techniques a. Sputter deposition Sputter deposition, also named cathodic sputtering, is based on the process of neutral atom release from a cathodic target, bombarded with positive ionized gas molecules which are accelerated by means of an electric field, and their subsequent -392- deposition on the substrate in order to form a thin solid film [3-5]. b. Reactive sputtering In this process a small quantity of the roactant in gas form is mixed to the inert sputtering gas and is introduced in the glow discharge or r.f. sputtering arrangement. This process is used to produce thin film of nitrides, carbides, oxides, hydrides, and sulfides [6]. c. Ion plating Ion plating is a combination of vacuum evaporation with radio-frequency sputtering considered as evaporation in a glow discharge or evaporation with a biased substrate. Similar to sputtering, a gas plasma discharge is set up between a cathode (the substrate) and an anode (the source of the material to be deposited) [3]. 1.4. Chemical Vapor Deposition Chemical vapor deposition (CVD) processes constitute a highly versatile and flexible means of applying coatings of most of the refractory metals and nonmetals [7,8]. CVD has become, within the last 25 years, the main technological method for producing films for semiconductor devices. The disadvantages, advantages, types of CVD and several other concepts will be discussed in more details in the following sections. 2- CHEMICAL VAPOR DEPOSITION Chemical Vapor Deposition (CVD) is usually defined as a process where gaseous precursors are introduced into a reactor and a stable solid film is formed on a substrate, by the reaction of chemicals -393- from the gaseous state making use of an activation energy [9-11]. In CVD process the formed solid film nucleates and grows on a substrate in an environment where a vapor phase chemical dissociation occurs [12]. CVD is a process where one or more gaseous species react on a solid surface and one of the products is a solid phase material. The several steps that must occur in every CVD reaction are as follows; - Transport of the reacting gaseous species to the surface, - Adsorption or chemisorption of the species on the surface, - Heterogeneous surface reaction catalyzed by the surface, - Desorption of gaseous reaction products, and - Transport of reaction products away from the surface [3]. The rate at which the initial state,low temperature reactant gases, changes to the final state, solid phase and product gases, depends on chemical kinetics and fluid dynamics. The properties of the deposited films and the deposition rate depend on the nature of the reactants and their purity, the amount of supplied energy, the substrate temperature, the ratio of the reactants, the gas flow rates, the system pressure, the geometry of deposition chamber, and the substrate surface preparation [2,3], Several types of reactions are classified under the CVD heading. These reaction types are pyrolsis, hydrogen reduction, halide disproportionation, -394- transfer reactions, and polymerization [13]. 3. ADVANTAGES Of CVD PROCESSES The main advantages of CVD technique consist in; - Possibly of forming various kinds of coatings, good quality deposits, high rates of depositions, - Possibility of making complex shaped bulky pieces, and - Producing uniform, reproducible and adherent layers of all classes of materials without defects and impurities in simple and cheap commercially available nonvacuum equipment [10,14]. CVD offers the following favorable aspects to semiconductor industry; - A wide range of silicon epitaxy, - A low cost, highly productive method of depositing, polysilicon, SijN( and high temperature SiOj and, - A proven technology for passivating silicon device by low temperature deposition of silicon- oxide and silicon-nitride. A wide range of metallic and inorganic coatings can be produced by CVD. In many instances CVD coating exhibit unique properties, which are difficult to produce by other methods [12]. CVD is an attractive method for applying ceramic coatings to composites because complex shapes can be coated without extensive equipment alternations. A wide variety of coating can be applied and coatings for unique properties can be produced [12] . -395- 4. DISADVANTAGES There are some disadvantages; among which are; - The use of commercially high temperature in many processes and chemical hazardus caused by toxic, explosive, flammable or corrosive gases [10]. - Often rather high temperature, Low pressures, - Heating of substrate itself sometimes difficult especially for large area, and - Accurate study of the reactor geometry needed prior to the coating of complex shaped components, etc.[14]. 5- DIFFERENT TYPES OF CVD CVD can be classified according to the type of their activation energy. Namely TACVD (thermally activated CVD), LCVD (Laser Assisted CVD), PACVD or PECVD (plasma activated or plasma enhanced CVD) , EBCVD (Electron beam induced CVD) and IBCVD (Ion beam CVD). Acronyms related to the process pressure are also in use; APCVD (Atmospheric Pressure CVD), LPCVD (Low Pressure CVD), and UHCVD (Ultra High Vacuum CVD). There also acronyms based on the precursors used. For instance, if a metal-organic compound is a precursor the process is called MOCVD, (Metal Organic CVD) [9,10]. 5.1. Thermally Activated CVD Thermally activated CVD uses thermal energy to produce a gas phase chemical reaction resulting in the formation of thin film on a substrate. Depending on the pressure value in the reaction environment, there are two classes of thermal CVD: - atmosphere-pressure CVD and low pressure CVD. -396- Both can be subdivided, in turn, into high- temperature and low temperature CVD, if the temperature is higher or lower than 500 °C, respectively [10]. 5.1.1. Normal-pressure thermally activated CVD In this technique gaseous reactants are di-luted by an inert gas (Hj, Nj, Ar or Helium) and flow over substrates maintained at high temperature where the total pressure is 1 atm. 5.1.2. Low pressure thermally activated CVD Is a process of thin film deposition on heated substrates at high or low temperature in a reactor under reduced pressure, usually in the range of 1 torr (0.01-100 torr), Fig. (2). Advantages - Supretation of autodoping from the substrate and vapor phase, - Improvement in film thickness and composition uniformity. - Control of the deposition rate only by the surface reaction rate, and - Decrease in defect number. Disadvantages - The lower deposition rates and the increased control maintenance [10]. 5.2. Plasna Enhanced CVD In this method, a glow discharge is produced in the gaseous reactant mixture which is maintained at a pressure of 0.1-1 torr under an RF plasma. -397- Fig.(1): Schematic representation of high temperature CVD reactors. (a)- horizontal plate reactor; (b)- barrel reactor; (c)- single wafer pedestal reactor; (d) pancake reactor; 1- quartz reactor; 2- cooling mantel; 3- RF coil; 4- substrate; 5- graphite susceptor; 6-pedestal; 7-exhaust. Fig. (2): Schematic representation of low pressure CVD reactors, (a): horizontal reactor; 1: gases; 2: gas flow controllers; 3: silica reactor; 4: three zone furnace; 5: wafer carrier; 6: wafers;7: pressure sensor;.8:filter; 9: pump; 10: furnace; 11: balast valve; 12: rotary vacuum pump 13: pump oil purifer; 14: exhaust, (b) vertical reactor.1: gas inectors; 2: wafer boats; 3: heaters; 4: vacuum pump. -399- Fig. (3). Advantages - The capability of producing conformal films at comparatively low temperatures (typically 200- 400°C), and - The possibility of obtaining films with amorphous structure on various heat-sensitive substrates. Disadvantages - Low deposition rate and efficiency, - Difficult control of film composition and thickness uniformity, - Inability to handle solid or liquid reactants, - Nonstoichiometric and inhomogeneous films, - Exposure of substrate and film to radiation damage, and - Complicated and expensive equipment and limited use on production scale [10]. 5.3. Photochemical CVD photochemical-assisted CVD, is based on the utilization of ultraviolet light to promote the decomposition of reactant gases. The reaction chamber having substrates are illuminated by an external UV lamp, Fig.(4). Advantages - Very low deposition (50-200 °C), - The occurrence of typical high-temp negative effects in semiconductor processing is minimized, and - The radiation device degradation is avoided. Disadvantages - Low deposition rate, -4(H)- rO-H VAC xuanC ana. rrm rrm-3| VAC Gas VAC (b) (O Fig. (3): Plasma-assisted CVD reactors; a: Laboratory vertical tube inductive reactor; b: production radial-flow planar reator; c: production hot tube reactor. 1: incoming gases; RF supply; 3: resistance heater; 4: graphite susceptors; 5: wafer; 6: RF coil; 7: wafer holder rotation; 8:vacuum pump system. Fig. (4): UV radiation-enhanced CVD system. 1: gases; 2:filters; 3: purge valve; 4: flowmeters; 5:HG resevoir; 6: mercury vaporizer; 7: vant valve; 8: reaction chamber; 9:substrate; 10: substrate heater {hot plate or IR lamps); 11: transparent window; 12: UV lamp; 13: throttle valve; 14:trap; 15: chemical pump; 16: exhaust diluent valve; 17: exhaust scrubber. - Nonuniform among wafers, and - The technique was applied to the deposition of amorphous silicon and insulating film. 5.4. Laser Induced CVD Laser induced CVD can be achieved using either pyrolytic or photolytic decomposition of gaseous phase molecules. Fig. (5) shows the experimental set-up for Laser CVD. Pyrolytic LCVD is based on local substrate heating by means of IR or visible laser light which is not absorbed by gaseous phase molecules. Photolytic CVD is based on electronic or vibrational excitation of the gaseous molecules using UV or IR laser. LCVD allows high rate one-step local deposition of insulating, semiconducting, and metallic materials, for production of microstructures. LCVD is potentially useful for the fabrication of integrated circuits [10] . 5-5. Electron Beam Assisted CVD This technique uses an electron beam to generate a plasma reaction in a small volume. Deposition occurring on a heated substrate (150- 500°C) located directly beneath that region. It allows high rate deposition of some dielectric films (SiO,-Si,N.) at low temperature (200-500 °C) [10]. 2 5.6. Ion Beam Assisted CVD In this technique a focused beam of ions is used to induce deposition. It is used in the repair of defects in photomasks by substituting deposited -403- -^E^fe Gas FIG. (5): Experimental set-up for laser CVD: 1: laser; 2: variable attenuator; 3: power meter; 4:lens; 5:reactor; 6:transparent; 7:pyrometer; 8:substrate; 9:localized deposited film; 10:gas supply; llrgas outlet. -404- carbon for missing chrome [10]. G. EQUIPMENT DESIGN Two different conventional systems can be used to produce CVD deposits, the open loop system and the closed loop system. Closed loop system were used for a small number of CVD processes especially for those with a low deposition rate, anyhow, this technique is not used any more. In open loop systems the reactant gases are continuously supplied and reaction product gases are continuously removed [14]. The equipment of thin film deposition using CVD generally contains the following units: - Gases or volatile liquid sources, - A gas distribution and mixing system, - A reaction chamber, - A system for providing the activation energy for the reaction and for heating the substrate, and - A neutralization system for the exhaust gases [9,10]. The design of a CVD system is affected by a number of factors; - The selection of the reactants, - Deposition temperature and pressure determine the materials that can be used in the construction of the system, - Moreover, the sensitivity of the process with respect to contamination from the gases and from air leakage influences the process economy (more sensitive gases and CVD systems), phase content, microstructure and the purity of the deposit [9]. -405- The equipment design depends on the type of activation energy, the initial aggregation state of the sources, the reactor operating principle, the substrate heating type, the reactor configuration, and the wall temperature [10]. The principle variables which require control and mointoiring are; pressure, temperature, reactant/product - activity/mass transfer and gas/ vapor - flow dynamics. 7. VARIABLES AFFECTING THE DEPOSITION RATE OF CVD In general, the variables affecting the deposition rate and film properties are; - the nature of the reactants and their purity, - the amount of energy supplied for decompose these gases, - the substrate temperature, - the ratio of reactants, - the gas flow rates, - the system pressure, - the geometry of the deposition chamber, and - the substrate surface preparation [10]. 7.1. The Effect of The Total Flow Rate In case of TiN deposits, the size of TiN crystal increases with increasing the total flow rate up to 700 cnrmin"1 but the size is constant when the total flow rate is grater than 700 cnrmin"' [15]. The carbon content of TiCxN„ coating layer is increased with increasing Cfy flow rate [16]. Mass transport limited rather than surface reactions limited. -406- 7.2. The Effect of The Deposition Temperature The size of the deposits, in case of TiN coating, increases with increasing the deposition temperature [15] • In case of SijN^-TiN composite the Ti content in the deposits decreased with increasing the deposition temperature and also the structure changed from amorphous to crystalline [17]. Increasing the deposition temperature for TiN films prepared by PACVD higher than 300 °C resulting in crystalline stoichiometric TiN while below 300 °C an amorphous TiN-Cl complex is formed [18] . 7.3. The Effect of The Partial Pressure The deposition rate is increased with an increase in the partial pressure of TiClj in case of TiN coating as long as the reaction is limited by mass transport [15]. 8- PREPARATION REQUIREMENTS 8.1. Safety in CVD All liquid chemicals are dangerous to some degree, therefore the following precautions should be taken into account in handling these chemicals. These precautions are; - Good ventilation should be provided in all areas where liquid chemicals are handled or stored, - Adequate equipment should be provided for the protection of face, hands, and body when working with solvents, strong acids and bases, - Use leak detection systems in all areas of CVD equipment, - Purge dangerous gas lines with separate N, cylinders, and -407- - Use auxiliary cut off valves in the line. 8.2. Substrates The general substrate requirements are as follows; - A thermal expansion coefficient similar to that of the film, - Chemical inertness to the growth conditions, - Chemical, mechanical, and thermal stability, and - A minimum number of surface defects. 8.3. Adhesion To have a good adhesive film the following requirement should be taken into consideration; - The mismatch in the thermal expansion coefficients between the substrate and the coating should be reduced by changing the deposition conditions or by predepositing a film which becomes an interface upon deposition of the final coating, - The homogeneous nucleation in the vapor should be eliminated by reducing the driving force, - Formation of brittle intermetallic compounds and pores should be avoided by predepositing an intermetallic layer, - Utilizing proper substrate cleaning procedures outside the deposition reactor to avoid surface contamination which affecting the adhesion [9). 9. APPLICATION OF CVD FILMS Application of CVD is extended to the manufacturing of all semiconductor deceives and integrated circuits. CVD technique has been used to; -4(18- - Produce power diodes and thyristors, Si, high speed transistor, Ge, electro-optic modulators, Gap, photo-diodes, GaAs, solar cells, CdTe, surface acoustic wave devices, A1N, GaN and ZnO for acousto- optic and electro-optic devices [19-21] . - Produce hetrojunctions, e.g., Ga/GaAs, SiC/Si, GaAo/Ge GalnPAs/GaAs, and AlGaAs-GaAs/GaAs [22-24]. - Produce insulators and dielectrics, e.g., SiOn, Si02:Si, Si,N4, A1203, BN, and A1N [25,26]. - Prepare Ni-Cr as resistors for hybrid IC's, Ni-Fe as magnetic films, TaAl for resistors for thermal print heads, and TaN and TaAIN as resistors for hybrid IC's [27-29] . - Improve the properties of cemented tungsten carbide tools with coating of TiN, Tic, or AL5O3 as a single layer or double-layer [15]. - Produce composite of SijN^-TiN for application in electronic devices [30] . - PACVD has been widely used for the production of Si3N4 passivation films and amorphous silicon films at temperature of 500 °C [31,32]. - TiN coatings is used as a corrosion resistance and in this case the pores parameters and the surface pretreatment must be carefully studied to minimize the existence of open pores in the coating [33]. TiC Coatings are used to improve surface properties of Tokmak elements, such as internal vessel walls, neutral beam injectors, divertors, and armour [33]. - TiN coating is coated on SUS 304 stainless steel to improve its corros.ion and abrasion resistivty to sea water and whirled sea sand [34] . - TiN films are used as heat mirrors for architectural applications and are very hard coatings [35]. - TiN coating is also used to improve the corrosion -409- resistance of the inner walls of stainless steel tubes in HC1 or HNO3 [36]. - PACVD is widely used for the production of amorphous and amorphous silicon carbide for solar cells which improved energy conversion efficiencies [37]. - PECVD is used for depositing TaSi2 on Si and SiO, [38]. - CVD is used to produce AI2O3 layers on Si to be used in the electronic industry in the fabrication of MOS field effect transistors [39]. - CVD is used for depositing tungsten silicide for integrated circuits manufacture [40]. PACVD is also used for in optical fiber preparation [41]. - CVD processes are used for depositing III-V compounds semiconductor materials [42]. - TiN coatings is used as a corrosion resistance and in this case the pores parameters and the surface pretreatment must be carefully studied to minimize the existence of open pores in the coating [43-48]. -410- aONCUUSIONS * CVD technique is used for forming various kinds of coatings, good quality deports, high rates of deposition, coating complex s ped bulky pieces, and the coated layers are uniform, reproducible and adherent free of defects and impurities. * A wide range of metallic and inorganic coatings can be produced by CVD. * CVD can be- classified according to the type of their activation energy. * The design of a CVD system is affected by the selection of the reactants, deposition temperature and pressure, and the sensitivity of the process with respect to contamination from the gases and from air leakage. * The equipment design depends on the type of activation energy, the initial aggregation state of the sources, the reactor operating principle, the substrate heating type, the reactor configuration and the wall temperature. * Application of CVD is extended to the manufacturing of all semiconductor devices and integrated circuits. -411- REFERENCES 1. M.S. Kim and J.S. John, Thin solid films, 107, 129-139, (1983). 2. K.L. Chopra, "Thin film phenomena", Mc Graw-Hill company, New york, London, panama, 10-69, (1979). 3. C.E. Morosanu, "Thin films by chemical vapor deposition"; Elseiver, Amsterdam, Oxford, New york, Tokyo, 31-61, (1990). 4. O.S. Heavens; "THin film physics"; printed in Great Britain by T.&A. Constable LTD Edinburgh, 23-56. 34. 5. J.C. Arderson, Editor, "The use of thin films in physical investigations", Acadimic press London and New York, 58-115, (1966) . 6. T. Hirai, and S. Hayaski, Material Science, 17, 1320-1328, (1982). 7. R. F Bunshah, J. G. Fish and J. M. Blocher, "Deposition technologies for coating and coatings, developments and applications", Noyes publications. New Jersey, USA, 1-362, (1982). 8. A. Kato, and Tramari, N., J. Crystal Growth; 29, 55-60, (1975). 9. D.S. Rickorby and A. Matthews, "Advanced Surface Coatings", ch. 7, Champman and Hall, New York, (1991). 10. C.E. Morosanu, "Thin Finis by Chemical Vapor Deposition", Elsevier, New York, (1990). 11. Jacek Korec, "Chemical Vapor Deposition Conference", 83-93, (1981). 12. M.G. Hocking, V. Vasantasree, and P.S. Sidkey, "Metallic and Ceramic Coatings", Longman Scientific Technical John Wiley fisons, Inc., New York, (1989). 13. Kasturi L. Chopra, "Thin Film Phenomena", 46-50, 1980, McGraw Hill Book Company New York. 14. E. Lang, "Coatings For High Temperature Applications", 33-60, Elsevier Applied Science -412- Puplishers LTD, England, (1986). 15. Moo Sung Kim and John S. Chun, Thin Solid Fims, 107, 129-139, (1983). 16. J. S. Cho, S.W. Nam, J. S. Chun, J. Materials Science, 17, 2495-2502 (1982). 17. Toshio Hirai, Shinsuke Hayaski, J. Materials Science, 17, 1320-1328, (1982). 18. M.R. Hilton, L.R. Narasimhan, S. Nakamur, M. Salmeron, and G.A. Somorjai, Thin Solid Films, 139, 247-260 (1986) . 19. Staff Article, Electronics, 45 (4), 41, (1972). 20. M.G. Craford, IEEE Trans. Electro. Dev., 29, 541, (1982). 21. J. Mimila-Arroyo, Y. Marfaing, G. Cohen-solal, and R. Triboulet, Solar Energy Mater., 1, 171, (1979). 22. S.M. Sze, "Physics of Semiconductor Devices", Wiley Interscience, New York, 2nd ed., (1981). 23. S.S. Chu, T.L. Chu and Y.X- Han, J. Appl. Phys., 60, 811, (1986) . 24. G. H. Olsen and H. Kressel, Electron. Lett., 15, 141, (1979). 25. K. L. Wang, S.A. Sailer, and P. Yang, IEEE Trans. Electro. Dev., 29, 541, (1982). 26. L. Bousse, N.F. De Pooij, and P. Bergveld, IEEE Trans. Electron. Dev., 30, 1263, (1983). 27. H. J. Belitz and O. F. Davie, U.S. pat. 2,742,691, (1965). 28. Staff Article, Electronics, 45(4), 41, (1972). 29. C.C. Lo and R.W. Kell, Hewlett-Packard J., 31(7), 22, (1980). 30. S.R. Kurtz,Thin Solid Films, 140(1986) 245-255. 31. R. Makake, S. Nakajima, and O. Takata, Thin Solid Films, 137(1986) L49-L50. 32. T. A. Maentyla, P.J. Helevirta, P.T. Siitonen, Thin Solid Films, 126, 275-281, (1985). 33.F. Brossa, J. Bohdansky, J. Roth, and A. P. -413- Martinelli, J. Nuclear Materials, 93&94, 474-478, (1980). 34. S. Motojima, and M. Kohno, Thin Solid Films, 137, 59-63, (1986). 35. S.R. Kurtz, and R.G. Gordon, Thin Solid Films, 140, 277-290, (1986). 36. J.S. Cho, Nam, S.W., and J.S. Chun, J. Material science, 17, 2495-2502, (1982). 37. H.Y. Kumagai 189-199 in proceeding of the ninth international conference on chemical vapor deposition 1984 The electrochemical society, incorporated. (1984) . 38. K.Hieber, M.Stolz and C.Wieczorek. 205-211 in proceeding of the ninth international conference on chemical vapor deposition 1984, The electrochemical society, incoporated, (1984). 39. Sung Woo Choi, Chul Kim, Jae Gon Kim and John Chun 233-257 in proceeding of the ninth international conferance on chemical vapor deposizion 1984 The electrochemical society, incoporated, (1984). 40- D.L. Borors, J.A. Fair, K.A. Monning and K.C. Saraswat 275-286 in proceeding of the ninth international conferance on chemical vapor deposizion 1984 The electrochemical society, incoporated, (1984). 41. P. Geittner 479-501 in proceeding of the ninth international conferance on chemical vapor deposizion 1984 The electrochemical society, incoporated, (1984). 42. Russeli D. Dupuis, 503-516 in proceeding of the ninth international conferance on chemical vapor deposizion 1984 The electrochemical society, incoporated, (1984). 43. J.S. Cho, S.W. Nam and J.S. Chun, J. Material Science, 17, 2495-2502, (1983). 44. C. Ernsberger, and J. Nicerson, J. Vac. science -414- technolog A, vol. 3 no. 6, 2302-2307, (1985). 45. R. Makaka, S. Nakayima, and 0. Tabata, Thin solid films, 137, L49-L50, (1986). 46. T.A. Maentylae, P.J. Helevirla, T.T. Lepistoe, and P.T. Siitonen, Thin solid films, 126, 275-281, (1985). 47. H. Itoh, K. Ato, and K. Sugioyama ,J. Materials Science, 21, 751-756, (1986). 48. R.L. Moore and Jr. Salvati, J. Vac. Sci. Technolog A, vol. 3 no. 6, 2426-2431, (1985). .415- First International Spring School & Symposium on Advances in Materials Science (SAAIS 94) 15-20 March 1994 EFFECT OF THERMAL BARRIER COATING ON LIFETIME OF FECrAlloy UNDER THERMALSHOCK CONDITIONS A- F. Wal-ieed and H. M. Soliman Atomic Energy Authority, Nuclear Research Center, Cairo, Egypt. Abstract : Thermal Barrier Coating (TBC's) in graded structure was deposited on FeCralloy. The ceramic part was partially stabilized zirconia, (ZrOi.YiOj) and the metallic part was Co- 29Cr- 6A1- 1Y. The graded coating consists of bond coatbond - ceramic graded zone / ceramic over coating. Two thermal cycling shock regimes have been conducted to evaluate the thermal barrier coating performance, and the effect of heating duration on the subsbrate/coating interdiffuction has been investigated. Scanning electron microscope, optical microscope and EDAX technique were used to discuss the results. KEYWORDS: Thermal Barrier, thermal shock, oxidation, zirconia. 1. INTRODUCTION: Metallic materials used at elevated temperatures for long term service are often strongly attacked by the environment (1-3). In most cases metal-gas reactions occur, leading to sulphidation. Such reaction can affect greatly mechanical properties of the structural materials (4-5). -417- FeCralloy is iron base alloy consists of more than 15% Cr and about 4.8% Al in addition to 1-3% Y as a reactive element. The alloys of Fe-Cr, Ni-Cr and Co-Cr exhibit the lower oxidation rates when chromium concentration is about 15% to 30% (6) . Therefore most of commercial compositions of these alloys contain about 16-25% Cr. Excessive amounts of Cr tend to promote embrittlement of in t.ermetall ic phase, another attribute of Cr,Oj is its slow growth rate (7). At elevated temperatures the presence of Al in the FeCralloy (<4.5%) forms A1203 scale, which significantly improves the protectiveness. Also, Al is very effective in improving cyclic oxidation resistance in air (8). The excellent oxidation resistance may exhibited in the presence of yttrium in the FeCralloys. The beneficial effects of rare metals (Y, Ce, ...etc) on the oxidation resistance of these alloys is well known for AI2O3 and C^C^ formers (9-13). One proposed mechanisms to explain their effects is the formation of oxide pegs, which mechanically key the scale to the substrate. A recent proposed mechanism (14) is based on the strong affinity of the rare earth elements (and another reactive elements) to sulfur, leading to strong sulfide formers which reduce the segeration of sulfur to the oxide scale-metal interface, thereby improving scale adherence to the substrate. Thermal barrier coatings (TBS's) consisting of plasma ceramic oxide and plasma sprayed metallic coat, in duplex or graded structure, are being developed for insulating air-cooled components in gas turbine and other heat engines (15-18). The bond coat material is an oxidation-resistant alloy of composition MCrAlY (M=Ni or / and Co). The aim of this work is to apply TBC' s on high temperature oxidation resistant alloys to improve the operating performance under severe conditions, and also to study the effect of thermal cycling on the interdiffusion between substrate and coat. -4 IS- 2. ErXJ^ERTMEEsTTAL : The ceramic part of the graded TBC's, is partially stabilized zirconia, Zr02- Y20j (93/7), agglomerated and sintered with particle 45 ± jiM. The metallic parts is Co- 29Cr- 6A1- 1Y. The graded coating consists of bond coatbond- ceramic graded zone/ ceramic over coat. These coatings, deposited on FeCrallby specimens in form of discs with 45 mm diameter and 3 mm thickness. The thermal barrier coated samples were thermally shocked in burner rigs. Two regimes of thermal cycling have been used in the first one the heating and cooling periods were 200 Sec, Fig. (1-a), while for the other the heating and cooling times were 300 and 200Sec, respectively. Fig.(1-b). Thermal cycle tests have been conducted in natural air. The surface temperature was measured using pyrometer, it was - 1000 and 1200 8C for the first and second regimes, respectively. The substrate temperature was measured by a thermocouple, it was 200 and 280 °C for the first and second group after cooling period. 3. RESULTS AND DISCUSSION: At temperature above 1000 °C the oxidation resistance of C^C^-forming alloys is markedly reduced due to transformation of the protective Cr203 scale to volatile Cr^Oj. To increase the operating temperature of the oxidation resistant of FeCralloy, it was coated with CoCrAlY and ZrO^-YjOj in gradecv. structure. The coat thickness was 300 |iM. The advantage of coating has been qualified in thermal shock at - 1000 °C and - 1200 °C using a high velocity burner rig. The microstructure of the graded TBC deposited on the oxidation resistant FeCralloy is shown in Fig. (2). Two groups of specimens have been tested, at 1000 and 1200 °C for the first and second regimes of thermal cycling mentioned before and shown in Fig. (1-a) and Fig. (1-b), respectively. The -4\9- (A) (J First regime V s0° — CJ 5 700 _^_ *\ ^ 6on - / I 500 o w 300 •Q \/ £ 200 J 00 i - 200 Sec -*+• 200 Sec —H Time Hsatmg poriod Cooling period (B) Second regime r 300 Sec *f»—200 Sec—H Time Hea ting period C oofing period Fig.(1 ): Typical thermal cycle of substrate. -420- .. . • .„*». 1.1 ... -• «r x .* #» Fig. (2): Microstructure of graded TBC on FeCralloy (X 200). -42!- shocked specimens were divided into four quarters, as explained in Fig. (3). Fig. (4) shows a specimen of the first group after shock test. The failure of these samples was observd after 50 cycles. The metallographic examination at the edge, Fig. (4-a), shows that the thickness is not much affected and the metallic part is slightly oxidized. However, few voids are observed in the coat-substrate interface. Maximum attack was observed at the middle of specimen where the burner strike the sample, as shown in Fig. (4- b). Therefore the coating thickness was affected in addition to voids formed at the coat-substrate interface. Fig. (4-C) shows an area close to that mentioned before (10 mm) apart from the center, oxidation of the metallic part of the coating was very clear, but the density of the voids at the interface was less. Interdiffusion between materials of different chemical composition can cause pronounced porosity near the interface because of kirkendall effect (19). This effect has been found in several MCrAlY coating-surface combinations (12, 20). . SEM photos for the same group are shown in Fig. (5). In the edge position, Fig. (5-a), few voids could be observed and no clear evidence of cracks in the over ceramic part. In the middle position where the burner strike the specimen, Fig. (5-b), shows the increase in the number of voids at the interface, and the cracks, in the over ceramic coating. These cracks, at 1000 C are mainly due to thermal fatigue and bond coating oxidation, according to Miller., et al (22). The authors reported that there was little phases change in plasma sprayed ZrOj-SY^Oj for temperatures below 1200 C. Therefore, there is no effect^ of phase transformation on failure of TBC at 1000 C. The zone between edge and middle was characterized by oxidation of the metallic part of the coat, as shown in Fig. (5-c). The second group of specimens were tested at 1200 C, with the second thermal cycling regime, Fig. (2-b). The edge zone, Fig. (6-a), shows crack -422- TBC_ | Substrate (a) before thermal shock Middle. Edge n «•• m m •» vrtm •*^^TB C Substate (b) alter thermal shock Fig.O): Schematics of the specimen. -423- lb* *> (A) Edge (B) Middle (CI Between middle and edge ^'^'K^e^inriSor"'-™ to •r**^-*»^;??.\C*jrt% Ti ^=J (A) Edge (B) Middle (C) Between middle and edge Fig (5): SEM for specimen of first group after thermal shock test (X 350). ;.:^v«^?.;v <~ •. ^ •. to (A) Edge (B) Middle (C) Between middle and edgt , , "••' ' 3£^?^J^1V5£«~» Fig. (7): SEM for specimen of second group after thermal schock test; middle zone. -427- in the ceramic part of the TBC, and oxidation of metallic part. Also, substrate voids are observed. Fig. (6-b) shows the middle zone, the thickness of the TBC is greatly affected, and in Fig. (6-c), a zone between middle and edge, the voids are less in size than those in the middle zone. Fig. (7), a SEM for the middle zone of the maximum temperature, shows the oxidation of the metallic part of the coat and the effect of thermal fatigue on the ceramic part. Tn this group of samples the number of cycles to failure was 12 cycles while it was 50 cycles for the first group. It can be concluded that the number of cycles to failure is affected by the time of heating cycle. At high temperature the life of the TBC is more dependent on the interdiffusion between coating and substrate (23, 24) . OONCMJLJ&XONS : 1. The high temperature potential of a coating is determined not only by its oxidation resistance but also by interdif fusions with the base materials. 2. Heating cycle duration, affects greatly the number of cycles to failure of thermal barrier coating. 3. Interdiffusion between bond coat and substrate causes the formation of voids at the interface. 4. The oxidation resistance of the metallic part of the graded thermal barrier coating affects its life time. /\CKIV»OWUE3X5MOSrr : The authors wish to thank the staff of IAW- KFA, Juelich, for providing the facilities to do this work. Also deep thanks to IAEA for the financial support of the fellowship. -428- REFERENCES r 1. Watson, J. W. and Levine, S. R., Thin Solid Films, 119, (1984), 185-193. 2. Stecura, S., Am. Ceramic Soc. Bull., 56 (12), (1977), 1082-1089. 3. Miller, R. A., Snoface and Coating Technology, 30 (1983). 1-11. 4. Kofstand. P., "High Temperature Oxidation of Metals", John Wiley, New York, (1966). 5. Kubschewiski, 0, and Hopkins, B. E. , " Oxidation of Metals and Alloys", Better Warths, London, (1967). 6. Stecura, S., Thin Solid Films, 73, (1980), 481- 489. 7. Birks, N. and Meer, G. H., "Introduction to High Temperature Oxidation of Metals", Ward Arnold Ltd. London, (1983). 8. Estrade-plato, D. J., and I. M. Haworth, C. W. Material Science and Technology, 2, (1986), 322. 9. Gulbramon, E. G. and Janson, S. A., in "heterogeneous Kinetics at Elevated Temperature", ed. G. R. Belton and W.L. Worrell, Plenum Press, (1970). 10. Smeggil, J. G., Mater. Sci. Eng., 87, (1977), 261. 11. Smialek, J. L., Metall. Trans. A. 18, (1987), 146. 12. Luthra, K. L. and Briont, C. L., Oxid. Met., 26, (1986), 26. 13. Huntz, A. M., Mater. Sci. Eng., 87, (1987), 251. 14. Ramanarganan, T. A., Ayer, R., Petkovic- Luton, R. and Leta, D. P., Oxid. Met., 29, (1988), 445-472. -42'J- 15. Miller, R. A., Levine, S. R. and Hodge, P. E., Proceedings of 4 International Symposium on Superalloys, ed. T. K. Tien, ASM, Metalspark, OH, (1980), 473-480. 16. Tajch man, P. and Bula, L., "Proceedings of 21'" International Thermal Spraying Conference, London, Vol. 7, (1989), 68. 17. Baharate Bhushan and B. K. Gupta, "Handbook of "Tribology" Materials, Coakings and Surface Treatments", Mc. Grow Hill Inc., New York, (1991). 18. Peterseim, J. and Jager, D. , "2n° Plasma-Technik Symposium", Lucerne, Switherland, Vol. 7, (1991), 241. 19. Peich, L, and Johner, G., J. Vac. Sci. Technol., A4 (6), (1986), 2582-2592. 20. Swindells, N., Mater. Sci. and Technol., 2, (1986), 250- 255. 21. Miller, R., Smialian, J. and Garlick, R., Advances in Ceramics-Science and Technology of zirconia. Vol. 3, ed. A. Heuor and L. Hobbs, Proceeding of the First Int. Conference on the Science and Tech. of Zirconia, Cleveland, (1980) . 22. Brennett, A., Mater. Sci. and Technol., 2, (1986), 257-261. 23. Grot, A. S. and Martyn, J. K., Ceramic Bull. 60, (1981), 807. 24. Nesbitt, J. A. and Gedwill, M. A., Int. Conf. Metallurgical Coatings, San Diago, (1987). -430- EGP5o^5e> First International Spring School & Symposium on Advances in Materials Science (SAMS 94) 15-20 March 1994. CHARACTERIZATION OF INDIUM-DOPED TIN OXIDE FILMS PREPARED BY SPRAY PY.ROLYSIS H.H. Afify, Department of Physics," Solid Slate Lab. National Research Center, Cairo, Egypt. M. Shafik Khalil, National Institute of Standards, Cairo, Egypt. R.S. Momlaz, Department of Physics and Mathematics, Faculty of Engineering, Suez Canal University, port Said, Egypt. Keywords: Doped tin oxide - Thin tilm - Spray Pyrolysis - Absorption edge - Energy gap - Optical - Optical and electrical Propeities. ABSTRACT Indium doped tin oxide films were prepared by pyrolyfic decomposition of a spray solution of chemically pure (MERCK) in CI3 . SnCI4. C2H5OH and water at different deposition temperatures and time. The transmission spectra of these films show a sharp decrease in transmission at the absorption edge and a good form of interference pattern. The film thickness as well as the optical constants n,k and a are calculated from the from the maxima and minma of the transmission spectra . Also the energy gap 4.15 eV was obtained. These values are plotted as a function of substrate temperature in 637~k to 798'k range. The sheet resistance of indium doped tin oxide films was. The obtained results are discussed in the light of the reported data in the literature. INTRODUCTION Transparent and conductive coatings of tin oxide have gained impurtancc in the field of research as well as technology because of their variety of applications. There is increasing usage of these films as thin film resistors, gas sensors [11. amirefleclion coatings [2] and transparent electrode liquid crystal display [3], optoelectronic devices and heteoiunction solar cells. Thin films of tin oxide can be prepared by a variety of methods.Among of these methods, reactive sputtering and evaporation, but they are quit expensive when large scale production is needed. Methods based on pyrolysis of an aerosol |5) represent a less expensive alternative. The setups used in the spraying methods have considerable differences from laboratory to another. This fact is being a possible source of the dispersion observed, in the reported properties of produced films by these methods. For practical puiposes this fact could be considered as an advantage, because each setup can be adapted and improved according to the applications of the produced films. Extensive work on tin oxide pure and doped with fuorine or antimony or both has been reported, doping with indium has not attract the researchers attention may be due to its great reduction to the conductivity of pure tin oxide. The high conductivity combined with high transparency for tin oxide films are the researchers goals. In the present work, spray pyrolysis technique is chosen for the deposition of tin oxide pure and doped indium thin films as it is simple and inexpensive process. The effect of substrate temperature and deposition lime on optical electrical properties of the deposited films has been studied to add more information to the literature in this respect. EXPERIMENTAL The pure tin oxide films were formed by spray pyrolysis technique using 0.6 Mol Sn CU in 98% ethanol. The tin oxide incorporated indium films were formed by adding 10 cm of 0.01 Mol CI3 soluble in distillled water to 100 cm3 of 0.6 Mol SnCIf in ethanol. The used home-made spray system was constructed in a chemical exhaust hood to avoide exposure to toxic vapors. The schematic drawing and other details have been given elsewhere |6]. The deposition temperatures were chosen between 673°k and 798V For this purpose, a stainless steel cylinderical block electrical furnace (5cm diameter and 15 cm height) has been used. During deposition, the temperature was automatically controlled within ± 2 °C by an electricronic regulator. The temperature uniformity over the entire substrate was necessaryto obtain homogeneous films. The temperature uniformity was checked by putting more than one chemically and ultrasonically cleaned glass substrate over the entire heater surface. The atomized fog of the spray solution was poured upon the substrates. It was found that the substrate at the central position of the heater surface had a uniform film. The spray time was varied from one to six minutes for each substrate temperature. Compressed filtered dry air was used as carrier gas with constant flow 5 L/min. The experiment was repeated more than one time to assure the. film reproduciblility. Optical transmission measurements were carried out in the wavelength range 0.2 to 2.0 u.m using shimadzu UV-Vis-Nir scaaning beam spectrophotometer model 3101s. This model is fully controlled by computer system. From the transmission interference pattern the refractive index n, the extinction coefficient k and the absorption coeffient* were calculated using a method reported in previous work [7J. The sheet resistance was measured by the conventional four-prope method. Film thickness was obtained from the experimentally traced envelope curves of the transmission spectrum TJ^and Tm values corresponding to the interference extremes in the very weak absorption region . The equations used to calculate the optical parameters n.k.a and the film thickness of weakly absorbing film from (T- X )curves are given elsewhere [9-12]. The transmission measurements are done in normal incidence mode. All spectra revealed interference fringes due to multiple reflection inside the film. -432- RESULTS AND DISCUSSION Optical Properties: The optical transmission of pure and doped indium tin oxide films deposited on glass substrates (2.0 x 1.2x0.1 cm) ismeasued at different deposition temperature ranging from 673'k to 798"k, constant deposition time (5 minutes ) and air flow rate 5 liter per minute was measured. Hg(1) shows the measured transmission curves for pu,e tin oxide films in 0.2 - 2.0 um wavelength range. It was observed that the decrease of transmission at the band edge is sharp for films deposited at temperatures above 698'k revealed interference maxima and minima due to multiple reflections inside the film> This could be attributed to the high crystallinity of the films deposited above 698>k and fine grained texture for films deposited below it. The fine and large grains were obsered orj these films by SEM and the degree of crystallinily below and above 698 k observed in x-ray diffraction had been reported elsewhere[8|. The effect of substrate temperature on transmission is the shift of peaks[1J with respect to each other . No significant change in optical transmission in the visible range 1-2 m the films deposited at temperature below 698 k. show a gradual decrease in transmission with the increase of substrate temperature. This may be due to the existence of free carriers with different concentrations. The decrease in transmission at the near infrared (1-2 um). for pure tin oxide and doped fluorine, was also observed by Gottlib el al [9]. The transmission curves for tin oxide doped indium with constant concentration at different substrate temperatures, constant deposition time of 5 minutes and air flow rate of 5 liters per minute are given in Fig.2. Small changes in opticl transmission in the wavelength range 0.2 (im - 2nm with an average value of 90 % has been observed. This may be due to the small diffcmces in film thickness. As shown in Fig.2, the significant feature of these curves is the sharp fall in transmission al the absorption edge as well the appearance of clear interference maxima and minima. This could be attributed to the high crystallinily and uniformity of the deposited lilms caused by indium doping. The transmission markedly increased in me wavelength rnge 0.9|im - 2.0«m as compared with that of puie tin oxide in Fig.3 shows a typical iransmiitance spectra of a 300 nm tin oxide doped indium film deposited at 773ck and 5 minutes depositioiuime. The peak maxima and minima envelops TM(>.) and Tm ().) respectively is shown in in Fig.3. Tj^and Tm in the curve are used to calculate the refraclive index n( /.) , extinction coefficient k and film thickness d. The calculated values of film thickness for the investigated tin oxide indium doped films as a function of substrate temperature and deposition time is shown in figures [4| and [5]. A comparable difference in thickness was observed below 723'k substrate temperature. The film thickness increased progressively with the deposition time at constant substrate temperature. The calculated values of film thickness are in the same order with that reported in the literature [10J. For (he determination of the fundamental absorption edge of pure and indium doped tin oxide films, which lies in the ultrviolet range, the lilms were deposited on quartz substrates. The absorption coefficient u = 4lk/;.of these films has been computed from Hie transmission dala. -4UH- Transmission (T%) 100 r 20O 400 600 800 1000 1200 1400 1600 1800 2000 Wavelength (nm) o 0 0 •798 K •773 K 723 K FigJ Transmission (T) of pure Tin Oxide thin films as a function of wavelengtM "X) anc at ceposition time of S minutes Transmi6sion(T%) 100 200 400 600 800 1000 1200 1400 1600 1800 2000 VVavelength(nm) o o o '693 K — 773 K 798 K ' quartz substrate Fig.2 Transmission curves of coped Tin Oxide tnin films as a function of wavelength O) and at a ceposition time 5 minutes -434- Transmission (T%) 100 200 400 600 800 1000 1200 1400 1600 1800 2000 Wavelength (nm) Fig.3 Transmittance(T) of Tin Oxide doped with Indium at temperature 773° K and deposition time 5 min. Thickness (nm) 350 670 680 690 700 710 720 730 7<0 750 760 770 780 790 600 Temperature CK) -*—1 mifuitt —•—J minuls* '—3 minute* -o— 4 mlnulct —X— Smlnule* Fig.4 Relation between substrate temperature (t) and thickness (d) lor doped Tin Oxide at different deposition times. -435- Thickness(nm) 3 4 Time (minutes) ~—637BK -•- 723°K -"— t*a°K -•-773 K -«- 796°K Fig.S Thickness growth as a function of deposition time for doped tin oxide at different temperatures 4 2 4.4 4.6 4.8 S S.2 S.4 5.6 S.8 Photon energy (electron Volt) -°—773 K -+-723 K —*-673 K Fig.6 Relation between square of absorption coefficient OC*and photon energy (e V) at constant aeposition time 5 minutes -43ft- The nature of the free carries transitions involved can bedeterimed on the basis of dependence of a on hv. For allowed direct transition, a is given by a = (hv - Eg)''where Eg is the separation between the minima and the maxima of conduction and valence band at the same wave number. Fig.6 shows the variation of aiwith hv, which is a straight line, indicating that the direct transition is dominat transition involved. The energy gap is the intersection point of the straight line with hv coordinate at at u = 0. The obtained energy gap values for and dopea tin oxide is nearly the same and equal to 4.15 ev\ The relation between refractive index for pure and doped indium tin oxide films and the covered wavelengths is depicted in Fig.7 and 8. The observed increase in n values in the shorter wavelength range can be associated with fundamental band gap absorption in these films. In large qavelength range (0.5 -2.0um) the n values has slight change with respect to each substrate temperature. The refractive index values (n) increases with increasing substrate temperature. It changes from 1.65 to 1.85 and 2.45 to 2.95 in the temperature range 673 - 798°k for the both pure and doped tin oxide respectively. These values of n are in good agreement with that reported in the literature [11]. The obtained refractive index values from 1.65 to 1.85 for pure tin oxide near the solar maxima of the solar spectrum (0.5 urn), are values larger than that of air smaller than that of silicon. It is therefore reduces the air silicon index mismatch when applied to the silicon surface, increasing the ability to capture solar radiation. In other words, the obtained small n values for pure tin oxide nominate it to be used as an antireflection for silicon solar cells. ELECTRICAL PROPERTIES Sheet resistance: The sheet resistance as a function of substrate temerature for pure tin oxide deposited at 5 minutes is plotted in Fig.9. The measured sheet resistance values for doped tin oxide films deposited on glass substrate at different temperatures and deposition times ranging from 673 - 798tk and 1-6 minutes respectively are represented in Fig. 10. Tin oxide doped indium have sheet resilance about three order of magnitude laiger than that of pure tin oxide. Airchison (12] found that the addition of atoms with a valence higher than four e.g Sb has a similar elfect to the reduction of Sn02 films and causes a significant increase in conductivity. Instead, the addition of trivalent atoms e.g In has the opposite effect, i.e it causes a decrease in conductivity. Accordingly the obtained results are in agreement with that of Airchison. From Fig.9 and 10, it is clear that the best deposition temperature and time to obtain films of pure and doped indium oxide with low sheet resistance are 748, 723^ and 5 minutes respectively. The sheet resistance decreases until it attains its minimum near 723 k . This behaviour may be due to the increase in both canier concentration and its mobility caused by large grain size. The increase in sheet resistivity beyond the temperature 723 k with the generation and increase of micro-pore size which cause the reduction in cross section of current flow. Also another possible reason of such behaviour could be the neutralization of charge carriers by the diffusion of alkali metal ions from the glass substrate, several reported data indicate such behavior of pure and doped SnOo with -437- refractive index(n) 4 — 200 400 600 S00 1000 1200 Wavelength (nm) -*- 798 *K -•- 773 *K ~*-723 *K Fig.7 Refractive index as a function of wavelength at oifferent substrate temperatures for pure Tin Oxide (oeposition time 5 min) refractive index(n) 6 5 •K 4 3 • t • PW' *"-»-£nir.: < 2 1 0 800 1000 1200 141 20 0 400 600 30 Wavelength (nm) •798 K •723 K •798 K Fig.8 Refractive index as a function of wavelength at different temperatures for doped Tin Oxice (deposition time 5 min) -438- Resistance (fl) 2500. 2000 • 150OI- 10001 5001- 0- - 660 680 700 720 7«0 760 760 800 Temperature (°K) Fig.9 Sheet resistance of pure Tin Oxid as a function of temperature at deposition time of 5 minutes Resistance (KG) lOOOOOir 660 680 700 720 740 760 780 800 Temperature (*K) -*" 3 minute* """"" 3 minvlta "•" * minutt* "*" S minute* Fig.10 Sheet resistance for coped Tin Oxide as afunction substrate temoer&ture at different deposition times -439- Ihe substrate Temperature [13-16] the obtained results here show nearly the same trend of that reported in literature. CONCLUSION Highly conductive pure tin oxide films with high visible transmiltance >85% have been prepared on glass substrates by spray pyrolysis technique. Tin oxide doped indium Ihin films sre also prepared. They had a high transmittance all over the covered wavelength range (rj.2jim-2.0nm) with relatively low conductivity . The substrate temperature has a high effect on the sheet resistance, but a slight effect either on the transmittance or energy gap of the prepared pure and doped flims. The results obtained are reproducible and are compared with those reported by other non vacuum methods employed to data The used technique is quite desirable for its simple approach coupled with its relevance in producing films of low resistivity and good optical transmission. REFERENCES [1] H.Pink, L. Treitinger and L. Vite, Japan J.Applied Physics, 19,513, 1980. [2] S.Sanz Maudes and T. Rddriguz, Thin Solid Films 69,183, 1980. [3] G.H. Heilmier. Sci. Amer. 22, 100. 1970. [4] K.L. Chopra, S. Major and D.K. Panadya, Thin Solid Films, 102,1,1983. [5] G. Blandent, M. Court and Y. Lagarade.-Thin Solid Films, 77,81.1981. (6] H.H. Afify, S.A, Nasser and S.E. Demian, J. Malerials Science\ Materials in Electronics. |7J H.H. Afify, R.S. Momtaz, W.A. Badawy and S.A. Nasser. J. Materials Science Wlalerials in electronics, 2,40,1991. [8] Under Publication [9] B. Gottlib, R. Koropecki, R. Arce, R. Crisalle and J. Ferron, Thin Solid Films, 199. 13,1991. [10] G. Sanon, R. Rup, and A. Mansingh, Phys. Stal. Sol (a), 128,19,1991. [11] H. Demiryonyt and K.E. Nietering solar Energy Materials 19, 79,1991. [12] R.E. Airchison, Australian J.Applied Sci., 5, 10, 195*. [13] Chitra Agashe, B. R. Maralhe, M.G. Takwale and V.G. Bhide, Thin Solid Films, 164-, 261, 1988. [15} A.L. Unaogu and C.E. Okeke, Solar Energy Materials 20, 29,1990. [16] G.Sanon, R. Rup, and A. Mansigh, Thin Solid Films 190, 281, 1991. .44(1- First International Spring School & Symposium on Advances in Materials Science (SAMS 94) 15-20 March 1994 On The Electrical and The Optical Properties of CdS Films Thermally Deposited By A Modified Source A.Ashour, N.El-Kadry and S.A.Mahmoud Physics department, Faculty of Science, Minia University, Minia, Egypt. CdS films were prepared by physical vacuum deposition technique on an amorphous (glass) substrate using modified evaporation source. Effect of processing conditions (the substrate temperature, the deposition rate and the film thickness) were considered in comparison with the normal evaporation source. Van der Pauw four-probe arrangement was used for measuring the electrical resistivity, which varies between 0.1x104 and 8.18x104 ohm cm. It decreases with increasing the film thickness, decreasing the deposition rate and the substrate temperature. This was attributed to the effect of crystallite size, degree of preferred orientation, internal microstrain and the stoichiometry. From spectrophotometric measuring, using approximate formula, the optical parameters (n, k, a) were evaluated. The values of the refractive indices and the energy gaps were evaluated and correlated with the preparation parameters. 1. Introduction In recent years there has been considerable interest in thin film semiconductors for use in solar cell devices and thin film transistors for flat panel displays[1-3]. Cadmium sulphide, CdS, belonging to the ll-VI group is one of the promising materials. The deposition of CdS films has become increasingly important due the widened industrial applications with a large number of uses(4- 6]. CdS has been the subject of intensive research because of its intermediate band gap, high absorption coefficient, reasonable conversion efficiency, stability and low cost[7,3]. The knowledge of the optical properties of CdS films is very important in many scientific, technological, and industrial applications in the field of optoelectronic devices, particularly solar cells. A broad variety of deposition techniques[8-10] permits the deposition of CdS films with desirable optical properties. Among these, spray pyrolysis, chemical bath deposition,.sputtering and vacuum evaporation are well established techniques. The vacuum evaporation technique is known to be suitable for preparation of CdS films for a wide range of applications [6.8.11]. In this work, the electrical and the optical properties of the CdS films, which deposited by a modified evaporation source under various preparation conditions will be investigated. The iilm resistivity as well as the optical constants n, k are calculated. The energy gap will be estimated from the optical absorption spectrum. -441- 2. Experimental CdS films were deposited by thermal evaporation from a resistivity heated quartz crucible containing high purity CdS powder and wrapped in a tantalum sheet. This arrangement was found to facilitate the loading of the CdS powder and to increase the lifetime of the heater. A small wad of quartz fiber wool was placed in the neck of the quartz crucible to prevent spattering of the CdS during outgassing and evaporation. The vacuum system (Coating unit E306A Edwards Co.) having a diffusion pump with a liquid nitrogen trap. Pressure of vacuum system during evaporation was kept at (1.3-2.0)X10"3 Pa outgassed at 150 ^C in the vacuum chamber prior to evaporation. The starting material was obtained from Fluka AG and of high purity (5N). Thickness and deposition rate were controlled by using thickness monitor and calibrated with a mechanical stylus method. The substrate used in the experiments were carefully cleaned microscope slides. The electrical resistivity of the prepared films were measured by the four - probe method of Van der Pauw[12] and the corresponding correction tables[13]. The potential difference and the current were measured by a Keithly electrometer (model 616). Four small indium contacts were deposited at room temperature on the CdS films which were arranged symmetrically around the circumference. Copper wire was attached to the contacts using silver paste. All leads were shielded to minimize induced currents and temperatures maintained at constant values across the sample to eliminate Seebeck voltages. The optical characteristics of the CdS films were measured with a Shimadzu spectrophotometer (model UV - 3101 PC). 3. Results and discussion The dark resistivity of the CdS films measured at room temperature is presented in table(1]. The resistivity values are ranging from 0.1x104 to 8.2x104 ohm cm for films with thickness 1000 and 1500nm, deposition rate of 0.8 to 2.2nm/s on substrate temperature of 180 to 220 °C. The data showed that, the resistivity increased as the substrate temperature and deposition rate increased, but it decreased with increasing film thickness. Similar behavior was observed by others[14,15]. For nearly the same specification ( thickness, rate, substrate temperature) of the films, the resistivity is higher in the case of the modified source than the normal one[14]. This found not to be due to the effect of crystallite size or strain[14,16] but seems to be to the better stoichiometry of the CdS film prepared by the modified source. -442- Tablep]: Values of electrical resistivity of CdS films at different deposition parameters. i t R Resistivity ; nTS (°C) (nm) (nm/s) (x104Ohmcm) J i 180 1000 0.8 0.40 { 200 1000 0.8 0.70 « 220 1000 0.8 5.50 | 200 1000 1.5 8.18 J 180 1000 2.2 5.20 i 1 i—.. 180 1500 0.8 0.10 | 180 1500 2.2 LOO ; 200 1500 2.2 3.00 { The variation of electrical properties as a function of the preparation parameter was highly correlated with structure characteristics of the prepared films[16]. Thus, the decrease in resistivity values as the thickness increase is a result of the increase of the crystallite size and the degree of preferred orientation of the films. On the other hand, the increase of resistivity with the increase of the deposition rate is due to the increase in the internal micro strain and decrease in both the crystallite size and degree of preferred orientation. Finally, the improve of stoichiometry as the substrate temperature rises, leads to increase of the resistivity. Thus, the dependence of resistivity is most probably the result of changes in the structure of the films. The optical transmission for CdS films is shown in Fig.[1] for samples prepared at different preparation parameters. It can be observed that, in general, the transmission is high. The refractive indices(n) were determined from the reflectance(R) data using[17]: R = [(n-1)2]/[(n+1)2] HI As shown in Fig.[2], the refractive indices of the films are greatly influenced by the deposition rate and substrate temperature. The refractive indices increase as the deposition rate increase, this is in agreement with Gottesman and Ferguson [17]. On the other hand, the indices decrease as the substrate temperature increases. The absorption coefficient, a, .and the extinction coefficient (absorption index), k, were obtained from the transmittance, T, and reflectance, R, using the approximate formula[18]. -443- T = (1-R)2 exp(-at) / [1-R2 exp(-2at)] [2] where t is the thickness and a is related to the extinction coefficient by a = 4nk/X [3) Fig.[3] shows the variation in k, as a function of the wavelength, which is found to be increased with decreasing deposition rate, and increasing substrate temperature in contrast as the behaviour of refractive index. Fig.[4] shows the wavelength dependence of the adsorption coefficient of CdS films. From this data the energy gap was estimated as shown in Fig.[5]. The curves revealed that absorption coefficient of these films is in the order of 104 cm'1 and decreases as the deposition rate increases and the substrate temperature decreases. This variation of the optical constant of the films can be correlated with their structural characterislic[19). The decrease of n and increase of k is due to the improvement of stoichiometry[20], increase in crystallite size(21] and decrease in internal strain[16]. The absorption coefficient, a, of CdS was found to follow the relational ,22], oE = A(E-Eg)0-5 [4] where A is a constant, E=hv is the photon energy and Eg is the band gap. The energy dependence of the absorption coefficient (Fig.(dj) indicates that these materials are of direct band gap. The optical energy gap is defined by extrapolating the linear portion of the absorption spectrum to E=0. Fig.(5] shows the plot of («hv)2 versus hv for different preparation parameters. The plots of CdS films show a straight line portion intercepting the energy axis at 2.23 and 2.3eV was found as a direct band gap material. Also, as shown in Fig.(5] at high thickness(t=1500nm) no effect of the rate and Tg on the energy gap, on other hand, at relatively small thickness, Tg and rate have no effect on the energy gap. Thus, using the modified evaporation source produces films of smaller energy gap than the normal evaporation one. However, it is worth to mention that the prepared film posses higher resistivity, indicating a change in the structure. 4. Conclusion The electrical and optical investigations prove that the used modified evaporation source(heated quartz tube covered with a small wad of quartz fiber wool) produces film of high resistivity and smaller energy gap than the normal -444- evaporation source {molybdenum boat). The results show that the resistivity decreases as the substrate temperature and deposition rate decrease. On other hand, it increases with increasing the thickness. The optical investigations based on the spectrophotometric characteristics have confirmed that the CdS films grown on glass show high transparency. The determined refractive index was smaller than that of bulk material. The deduced extinction coefficient was relatively smaller than those reported in the literature and increased sharply at the bandgap region. The films exhibited direct transition in the range 2.23-2.30eV. This band gap values are smaller than those obtained by other thin film techniques. No difference is expected on the £„ for the examined thickness. This can be assumed as a continuous film.S o surface effect are negligible. Acknowledgment The authors would like to thank Prof.Dr.A.A.Ramadan, Physics Department, Minia University, for his help and useful discussions. 5. REFERENCES [I] K.L.Chopra and I.I.Kaur, in: Thin Film Device Application, eds.K.LChopra and T. C. God (Plenum Press, New Yourk, 1983) [21 S.Chandra, R.K.Pandley and R.C.Agarwal, J.Phys.D:13,1757,1980. [3] K.LChopra and S.R.Das, Thin Film Solar Cells (Plenum Press. New York, 1983). [4] R.Hill in: Thin Solid Film Active and Passive Devices, ed. T.J.Coutts (Academic Press, New York, 1987)p.487. {5] M.Ueno, H.Minoura. T.Nishikavva and M.Tsuiki, J.Etektrochem.Soc, 130,43, 1983. [6] J.Santamaria, I.Martil, E.lborra, G.Gonzalez Diaz and F.Sanchez Queada, SoI.Energy Mater., 28,31,1990. [7] F.EI Akkad and M.Abdel Naby, SoI.Energy Mater, 18, 151, 1989. [8] S.N.Sahu and S.Chandra, Sol.Cells, 22,163,1987. 19] LP.Desmukh, A.B.Palwe and V.S.Sawant, Sol Energy Mater, 20,341,1990. [10] A.Bennouna, E.LAmeziane, A.Haouni, N.Ghermani, M.Azizan and M.Brunei, SoI.Energy Mater., 20,405,1990. [II] J.W.Orton, B.J.Goldsmith, J.A.Chapman and MAPowell, J.Appl.Phys., 53, 1602,1982. [12] L.J.Van der Pauw, Phil.Res.Rep., 13,1,1958. [13] A.A.Ramadan, R.D.Gould and A.Ashour, Thin Solid Films, 239, 272,1994. [14] A.A.Ramarian, R.D.Gould and A.Ashour, Int.J.EIectronics, 73, 717, 1992. [15] S.Ray, R.Banerjee and A.K.Barua, Jap.J.Appl., 19,1889, 1980. [16] A. Ashour and S.A.Mahmoud, 1§J Int.Spring School & Symposium on Advances in Materials Science(SAMS 94)15-20 March 1994,Cairo.Egypt. [17] J.Gottesman and W.F.C.Ferguson, J.Opt.Soc.Am„ 44, 368,1954. -445- [18] I.Martil Oe La Plaza, G.Gonzalez-Diaz, F.Sanchez-Quesada and M.Rodriguez-Vidal, Thin Solid Rims, 20.31.1984. [19] A.A.Ramadan, R.D.Gould and A.Ashour, to be published. [20] K.Yamaguchi. N.Nakayama, H.Matsumoto, and S.lkegami, Jap. J. AppI. Phys., 16.1203.1977. [21] F.Tepehan and N.Ozer, Solar Energy Materials and Solar Cells, 30,353, 1993. [22] O.P.Agnihotri and B.K.Gupta, JapJAppl. Phys., 18,317,1979. -446- 100 ,.. r> A /V \ / A ^ so 60- i80°C 0-8 nm/s — 200 °C 0-8 nm/s 200 °C 1-5 nm/s-Tf-s- AO- 220°C 0-8 nm/s.»-— (a) 20 100 11 -i i '.,w\ 60- / «/vt( Tv 60- 180 C 0-8 nm/s 180 °C 2-2 nm/s 200°C 2-2 nm/s (b) 700 800 900 Wavelength( nm) Fig.[1]: Optical transmittance of CdS films at different preparation conditions. (a) t=1000nm, (b) t*1500njn .447- "3 — 200 C l-5nm/s -200 C 0-8nm/s 2J 0-5 (c) 3-5 -~180C 2-2nm/s — I80'c 0-flnm/s — 200*C 2-2nm/s (b) 0-5 SCO 5 00 700 800 SCO Wcveiength(nm) Fie.[2]: Wavelength dependence of refractive index of * cds fiias at different preparation conditions. (a) t-lOOOnm, (b) t»lSOOnm -448- -c-220 C 0-8 "m/s — 2C0°C 0-3 nm/s — 200 C 1 -5 nm/s 500 600 700 800 SCO Wavelength(nm) rig-. [31: Absorption index characteristics of CdS film* at different preparation conditions. (a) talOOOna, (b) t-lSOOnm -449- 5. e' 220 C 0-8 nm/s 200 °C 0-8 nm/s 200*C 1.5 nm/s 1- t« 0 (a) x 180 C 0-8 nm/s 200 °C 2-2 nm/s 180°C 2-2 nm/s j- 500 S00 7C0 8C0 SCO IOCO Wavelength(nm) Fia [41 : Absorption coefficient characteristics of CdS filns at different preparation conditions, (a) t-lOOOmn, (b) t-lSOOnm -4M»- 220 C 0-8"m/s 200 *C 0-8nm/s 200°C 1.5 nm/s CO ieo'c 0-8nrn/s — 200*C 2- 2 nm/s (b) 2-6 2-8 Energy (eV) Fig.[5]: «XE)2 voraua E ploc of CdS filma at differene preparation conditions, (a) fcalQOOnm, (b) t»150Own y>45V First International Spring School & Symposium on Advances in Materials Science (SAMS 94) 15-20 March 1994 Substrate Temperature Effects on Tin Oxide Films Prepared by Spray Pyrolysis IULAfify, F.S.Terra** and R.S.Momtaz* Solid State Physics Department, National Research Centre, Cairo, Egypt * Department of Physics and Mathematics, Faculty of Engineering, Suez Canal University, Port Said, Egypt. Key words : Spray pyrolysis-tin oxide doped with indium-x-ray diffraction- SEM-iexturc coefficient-sheet resistance-substrate temperature. Indium-doped tin oxide films were prepared by spray pyrolysis technique at different substrate temperatures ranging from 4(K)-525*C. Texture coefficients for (2U0) and (112) reflections of tetragonal SnO, were calculated. The surface morphology of the prepared films was revealed by using scanning electron microscope. A denderile structure was ol)served in the films deposited at substrate temperature = 525*C . The obtained sheet resistance results were correlated with those obtained from x-ray diffraction analysis and scanning electron microscope. I. Introduction Economical and stable transparent conductive SnO- films are of considerable interest due to their application in solar cells, optoelectronic devices, thin film resistors, antircflection coatings, photochemical devices, and electrically conductive glass. It has been found that the spray technique is most suitable for obtaining tin oxide films in a large area sulwtrate applications. The surface to|x>graphy of undoped and doped SnO- films have been studied by scanning electron microscopy "' "' (SEM). Tin oxide films not only have good adherence but also high chemical stability on a variety of substrates. The structural and electrical properties of pure and doped SnOj with As, F and Sb have been investigated as a function of substrate temperature* "4* "'• \ Inadequate information on indium-doped tin oxide films reported in the lileratuie. In this work, x-ray -453- diffraction, surface topography by SEM, and sheet resistance for Sn02 : In films as u function of substrate temperature will be studied. II. Experimental An inhousc spray pyrolysis set-up was used to prepare the thin films of indium-doped tin oxide . The structural details and schematic representation of the spray system arc given elsewhere. The spray solution is composed of 0.6M Sn Cl4 (Merck), dissolved in 98% ethyl alcohol. In Cl3 (Merck) (10 ml of 0.(11M ) in distilled water was added to 100 ml of 0.6 M Sn Cl4 in ethyl alcohol, in order to obtain indium-doped tin oxide films. Chemically and ullrasonically cleaned glass substrates (2.2 x 1.1 x 0.1 cni"*) were used. The substrate was placed on the surface of a cylindrical stainless- steel block (diameter = 5 cm and height = 15 cm), electrically heated to attain the required substrate temperature, the temperature was measured by a Pl,Pt-Rh thermocouple with a temperature monitor Philips thermostat IT 22K2A, which measures the temperatures from -lOO'C- 1000'C. The furnace temperature is controlled by a proportional temperature controller Eurolherm to + 2*C. A fully automatic ac-voltage regulator JESPEC model ST 3000 was used to stabilize the power, which supplies the spray system. Compressed air was used as a carrier gas, with a flow rate of 5L/ min. The spray time was kept constant at 5 min. for the whole substrate temperature range 400-525'C. The effect of film thickness on the properties of the deposited films was previously studied. X-ray diffraction analysis was carried out using Philips PW1710 x-ray dilfraclomeler, with CuKa radiation X = 0.1542nm) and Ni filler at 40 kv and 30 inA. The sheet resistance was measured by conventional four-probe method using 5 digits ammeter and voltmeter (G.L.A. Electronica). (SEM) Siemens, was used to study the surface morphology of the deposited films. III. Results and Discussion Ill.l X- ray Diffraction Analysis The x-ray diffraction tracings of as-deposited indium doped tin oxide films at different substrate temperature ranging from 400-525'C are presented in Fig. (I). Interplaner spacings and their relative peak intensities are calculated and compared with the standard values given in the ASTM cards for SnO- powder. The d-values with their relative intensities together with the ASTM standard values are given in Table (1). The deposited tin oxide films have noncrystalline tetragonal rutile structure with preferential orientation along [2(XI] at all the substrate temperature range. Traces of cubic SuC^, are delected at substrate temperatures 425-525*C, by the presence of (III) and (400) reflections. Also a reflection ofd= 2.05 "A appeared, which may be due to the presence of SnO in the films. -454- Table (1) ASTM Cards Substrate Temperature ("O 400"C 425*C 450*C 475'C 500*C 52S"C Material Syilim SnOj Cubic 2.842 100 lit -a« mm* 2.S31 32 2.831 6 2.S49 4 2.831 5 2.831 8.6 SnOj Tciiagonal 1369 24 200 2.363 100 2.363 100 2.363 100 2.363 100 2.363 ICO 2.356 JCO SnO 2.03 2.05S 21 2.0535 22 2.050 23 2.053 22 2.053 20 2.049 24 Sn02 Tcttagoiul 1.439 17 112 1.455 80 1.451 50 1.451 50 1.451 64 1.449 75 1.443 62 Sn02 Cubic 1.2Z6 S 400 1.236 17 1.233 14 1.230 15 1.233 18 1.236 16 1.230 17 Sa02 Tctilgonal 1.1840 3 403 I.18S 6 1.182 6 1.1S0 6 1.1S2 6 I.1S0 6 1.177 6 LA. Jl. ** . LJU- 80 76 7; 08 Wl 00 ib Ji 48 44 to 36 j'f^ 79 «g. C) X-ray di'ffnclimi (races- ofSuOj films al diffcrciil xiihslraic Icmpcralo -456- The texture coefficient TC (hkl) is calculated from the x-ray diffraction results us follows <9',t,):- I ( hkl ) /1. ( hkl ) TC(hkl) = l/N£Nl(hkl)/l.(likl) where, I, is the measured x-ray diffraction peak intensity, !• is the corresponding intensity, given in the ASTM cards, for the powder, N is the lunnoer of reflections. The texture coefficients arc calculated for the reflections (2(H)), (400) and (112) of the tetragonal tin oxide (SnO.) doped with indium at the covered substrate !ciii|K*rature range . Variations of the texture coefficients of (200) and (112) reflections with the sutislrale temperature arc shown in Fig. (2). TC (400) has the same behaviour as TC (2(H)). hut with smaller peak intensity. The growth mctfianisru could be explained from the standard deviation, u , of the x-ray intensities of the reflections, as follows C«'°): OS \f •• • . M wliere 1 is (be relative intensity of the (hkl) plane. The dependence of the standard deviation, a on the suhstrate temperature is shown in Fig. (2). The results are in good agreement with those reported by Chitra Wu) el ai. Tlic decrease of a at lower sulistrate (cniperatiires followed l»y stahility at higher ones, may he due to the homogeneous nuclcation and preferential SnO, growth for the deposited films. At substrate temperature of 400'C, the comparatively larger values of o may lie caused hy noiiluiniogencous Init preferential SnOj growth, which is confirmed hy the d-vahies. From the texture coefficient in Fig. (2) it appears that at the sultslrale temperature range 400-45O*C, the preferential [2(H)] orientation is enhanced , while [112] orientation is suppressed . However, at the suhstrate temperature range 450-500*C the preferential [2(H)] orientation is suppressed while the [112] orientation is enhanced. At 525*C the [20(1] preferential orientation is comparatively larger than the [112] orientation. II 1.2. Scanning Electron Microscopy (SKM) Details of (he surface morphology of as- deposited indium-doped (in oxide films al the substrate temperature range 400-525'C, were studied hy SUM. Tin: SUM micrographs are presented in Fig. (3) (A.I3.C,ID.ri,F). -457- nm TNckma, (nm) SutiiUtit T«mp»r»lui».T, (*C) Film Tni 3 f I ^ too « SuSil/jtr T.mp«/ltu»«, 7, (•{) -458- Kig. (3) SEM micrographs at sul)Slralc temperatures 4(X)*C, 425*C, 450*C, 47.VC. 5IK)*C and 525*C, as illustrated in (A), (B), (Q, (D), (E) and (F), respectively. -459- D \&&ma&msi *a*5 -460- The observed rounded grains could he attributed to the non-uniformity of the deposited Sn oxide, layer caused hy the broad droplet size Ihul forms " island structure" of different sizes. The difference in islands size leads to a competitive growth process and thereafter, columns with domed lojis are developed. At substrate temperatures > 500"C, the solvent vaporizes as the droplet approaches the substrate, then the solid melts and vaporizes. The vapour diffuses to the substrate to undergo a heterogeneous read ion " '. Consequently, the growing islands will have approximately the same size. In this case the boundary of the island extends and no growth-competition occurs. This leads to columnar structure of the film, in which the columns grow perpendicular to the substrate plane. This was assumption reported by Fan el al. » '. The variation of the sheet resistance of the prepared as-deposited indium-doped tin oxide film with substrate temperature is shown in Fig. (4). The sheet resistance decreases with substrate temperature, giving a minimum value at 45()*C then it increases again. This agrees with the results given previously W \ Correlating the results of the sheet resistance and the texture coefficient, one can ol)serve, that at substrate temperatures range 450-500*0 the sheet resistance decreases due to the enhancement of [200] preferential orientation of the grown film, while the increase in the sheet resistance from substrate temperature 45()'C to 50O*C is referred to the suppressed preferential [200] orientation and the increase of [I12] orientation. At substrate !cin|>crnturc 525*C the [200] and [112] orientations are present with nearly the same degree. In conclusion, comparing this result wiilt the SUM result, the denderile structure leads to an increase in the sheet resistance. From the standard deviation, a irrespective of the change of the sheet resistance with substrate temperature, preferential Sn02 format ion is confirmed to be present al the covered substrate temperature range in (lie present work. References 1. E.Shanthi, V.Dutta, A.Iiancrjce and K.LChopra, J. Appl. J'hys., 51,12(1980) 6243. 2. E.Shanthi, A.Bancrjee, V. Dutlsi and .L Chopra, J.Appl. I'hys. 53,3 (1982) IK 15. 3. D.Bclangcr, J.P. Dodelel. B.A.Lnmlxis and J.L. Dickson, J. Electroclicm. Soc. 132 (1985) 1398. 4. K.Suzuki, N. I lusimoio, T.Oyaina, J. Shiiiiizu, Y. Akaoand H.Kojima, Thin Solid Films, 226(1993) KM. -461- 5.11.Kancko and J.Miyake, J. Appl. Pliys. 53 (19X2) 3629. 6. K.I I. Yoon and J.S.Song, Solar Materials and Solar Cells, 2K (1993) 317. 7. G.Sanon.R.Rup and A.Mansingh, I'liys. Slal. Sol. (a), (1991) 109. H. Ilidcya lida, Toshio Mishuku, Atsuo, ho, Knumci Kalo. Mislsuyuki, Yanianka and Yulaka Hayashi, Solar Energy Materials, 17(I9XX)407. 9. Chilra Agashe, O.K. Maralhe. M.G. Takcwalcand V.G. Bhidc, Thin Solid Films, 164 (1988)261. 10. Chilra Agashe. M.G.Takewale :nul V.G. Uhkle, Solar Energy Materials. 17(19X8) 99. ] I K.li.Sumbrtfcini and G.K. Bliagavol. Tliiit Solid Mints. 78 (1981) 35. 12. A.Czapla, ILKiisior and M.Bucko.Thm Solid Films, 182 (1989) 15. 13. J.P. Hirlh. K.L. Moazed, G.Hass and R.E.Thun, Physics or Tliin Films, Vol.4, Academic Press. New York, (1967) Chap. 2, p. >JH. 14. J.C.Vigvie and J.Spitz, J.EIecirochem. Soc: Solid Stale Science and Technology, 122 (1975)585. 15. Y.D.Rm, X. P.Li. J.Yong and J.P.U. Phys. Slal. Sol. (a) 134 (1992) 157. -462- First International Spring School & Symposium in Advances in Material Science (SAMS 94) Cairo, 15 - 20 March 1994 EFFECT OF TEMPERATURE ON Sn02 THIN FILMS PYROLYTICALLY DEPOSITED R.S.Momtaz*, S. Darwish^, H.H.Afify*, Department of Phytic* and Mathematics), Faculty of Engineering, Suez Canal University, Port Soid, Egypt. Department of Physics, Faculty of Science, Minia University, Minia, Egypt. Department of Physics, Solid State Lab. National Research Center, Cairo, Egypt. Abstract Pure and incorporated Sn02 thin films with 8% Cadmium are prepared at substrate temperatures ranging from 400 to 525"C. The produced hard, Transparent and conductive thin films are prepared using spray pyrolysis technique. The film thickness, transmittance percentage, refractive index, figure of merit and energy gap values are investigated, the conductivity as well-as x-ray diffraction -nalysis are carried out. The effect of substrate temperature on the previously measured parameters either ir. pure or doped samples and the band gap are also examined. It is found tht the existence of cadimum through the matrix of SnC>2 influences the preferred orientation. 1. Introduction Highly transparent, hard, conductive oxides such as Sn02 represents a promising materials. This is due to their suitable properties for different applications. Tin oxide was used successfully in solar cell fabrication^, liquid crystal displays^, optoelectronic devices^, photoelectrochemical cells^3"°J, thin film resistors^4'5!, antireflection coatings and -463- . cells^3"81, thin film resistors^4,5^ antireflection coatings and gas sensors'7^ The spray pyrolysis technique is one of the popular techniques for obtaining such films. This is because it is simple, inexpensive, fast, effective in mass production and does not require evacuation. Many investigations have been performed to determine the characteristics of the produced Sn02 films. The previous trials were devoted towards the production of highly transparent and conductive thin films. So, doping the Sn02 matrix with different elements such as halogens^9"17* or trivalent and pentavalent elements are, required for such investigation as well as the optimization of the film production conditions. In this work, cadmium has been chosen to develop the properties of Sn02 films. Substrate temperature has been taken as a varying parameter at a constant concentration of 8% cadmium. The effect of doping at different temperatures on the optical properties (transmittance (T%), refractive index (n), thickness (t), energy gap}, electrical conductivity and the structure of the film are carried out. 2. Exerimental Pure and 8% cadmium doped Sn02 thin films were prepared on glass and fused silca substrates using the spray pyrolysis technique as described elsewhere^9'. The pyrolysis reactions are carried out at different temperatures ranging from 400 to 525'C according to the equation: m c SnCI4 + 2H20—* -™' » Sn02 + 4HC1 A stainless steel furnace (5 cm diam. and 15 cm height) has been used in combination with an automatic electronic regulator to adjust the test temperature within ± 2"C. A transparent, glass like, clear thin film is produced as a result of spraying a solution of 0.6 M SnCl4 in ethyl alcohoj (98%) and solution of 0.6 M CdCl2 in water. The rate of driving gas (compressed filtered air) was 5 1/min. -464- Ultrasonically cleaned substrates of the dimension of (2.2x1.2x0.1) cm. were used. For each substrate temperature, a run was carried out for deposition times from two minutes up to ten minutes. Optical spectra were taken using a Varian DMS 100 UV/VIS double beam spectrophotometer. Through the range of 200-900 nm. From the produced interference patterns, the refractive index n, the thickness of the films, and the energy gap were calculated^1**-19!. The thickness of the films were obtained from TM - Tm envolope values for each (T-?0 interference extrema in the very week absorption region. The thicknesses used in this study ranged between 60 to 400nm. The sheet resistances were measured at room temperature by the conventional four-probe method using a digital multimeter through a 1 cm2 fixed area. Structural analysis was carried out using a JOEL diflfractomatcr (model JDX-60 PA) with a Ni filter and Cu Ka radiation (k - 0.1542 nm) at 35 kV and 15 mA. 3. Results and Discussion 3.1 Optical Properties 3.1.1 Effect of substrate temperature and doping on transmission Substrate temperature was varied from 400 to 525°C. Figure (1) illustrates the transmittance spectra for the investigated samples of doped tin oxide with 8% Cd. It is observed that the transmission percentage decreases by increasing the substrate temperature. Figure (2) represents (T-X) for thin films of pure SnOo and doped 8% cadmium of the same thickness formed at the same substrate temperature. The transmission is decreased by = 10% for doped film which exhibit a crystallinity higher than the pure one. This is reflected as a maxima in the interference due to the multiple reflections inside the films. -465- 100 80 • /t Si 60 h i ('/ 40 20 450. C —r- 500 C 5JS*C 20O 3^0 480 620 7 60 900 A {nm ) Fig. (1) Effect of substrate temperature on Transmission (T%) of doped Sn02 with 8% Cd as a function of wavelength (X). -466- 100 80 GO 40 20 pure SnO^ doped Cd 200 340 480 620 760 900 K ( n m ) Fig. (2) Transmission (T%) of pure Sn02 and 8% Cd-doped Sn02 thin films at 450°C as a function of wavelength (I). -467- To study the effect of reflection on transmission values, the optical measurements were carried out by a double beam spectrophotometer. The transmission of films was measured by ratio-recording technique^21^ which measures directly the ratio of transmissions of films of different thicknesses. A thin sample was placed in the reference beam, while the thicker one was placed in the sample beam. This leads to a measurement of the transmission of thin film without reflectance losses. Figure (3) represents the cases of using clean glass substrate as a reference and a very thin film (= 30 seconds depossition time) as a reference for the measuring of the same sample. It is clear that the effect of reflection is increased by about 5% of the transmission values all over the range of measuremimt. We used the thin film sample as a referance in all measurments. 3.2.1 Refractive Index (n) The optical parameters, the refractive index, n, and thickness of the thin films, t, were calculated from the interference extrema of the T-X transmittance spectra. The method of calculation is described elsewhere^18,19! and is based on the analysis of the peak maxima TM (1) and minima Tm (1) envelops as shown in Fig. (4). Figures (5) and (6) represent the values of the refractive index, n, as a function o£ wavelength, X, at different substrate temperature for pure and Cd-doped Sn02 films respectively. It is clear from the figures, that the values of n versus X generally decrease with increasing the substrate temperaure. The main feature of these curves is the existence of two regions. First, the nearly steady region at X > 520 nm. The second, is the fluctuating region existing near the ultraviolet spectra for pure and Cd-doped Sn02. The fluctuation in the refractive index near the ultraviolet region may result from the existence of multi-oriented Sn02 macrocrystalline, each orientation having a different refractive index. The x-ray diffraction results for these films -468- 9 00 Fig. (3) (T-l) curves for a sample with a glass as reference and with thin film 30 s as reference. -469- 100 r -2 900 Fig, (4) A demonstration envelope curve to show how to get the values of refractive index (n). -470- show a large improvement in the crystallinity with increasing substrate temperature. For the films deposited at high temperature, there is one main preferred orientation (200) having a single value of refractive index. The obtained refractive index values are in good agreement with that reported by Demiryont and Metering'22^ on Sn02 prepared by evaporated and spray pyrolysis techniques. In general, the lower values of the refractive index (== 1.8) near the solar maximum (500 nm) for pure and Cd-doped Sn09 nominate these films for use as an anti-reflector for solar cell'-1'. 3.1.3 Effect of doping on the energy gap The ratio of transmission of two films of different thicknesses, which is obtained directly from the ratio-recording technique, is given by T^ a exp (a A t), where T1>2 is the ratio of transmission, A t, is the thickness difference of the two films. We can accurately calculate a A t without determining A t and the above equation can be written as 1/2 aAt«a0(hu-AE) By plotting (a A t)2 against h^, we get a straight line with a slope of OQ, the abscissa intercept giving the energy band gap AE, this is shown in Pig. (7). It is seen that three curves correspod to different thicknesses for pure and doped films. Each set has a common interception, giving a value of 4.05 eV for the doped films, and a value of 4.1 eV for pure 9 Sn09, the same as recieved before' '. These data show that the effect of doping by 8% cadmium on the value of the energy gap is very weak. -471- 300 400 500 600 700 800 900 X (nm) Fig. (5) Refractive index (n) as a function of wavelength at different substrate temperature for pure Sn02. -472- N 425 C • 450 *C » 500 *C 0 525'C 300 4.00 500 600 700 800 900 \ (nm) Fig. (6) Refractive index (n) as a function of wavelength at different substrate temperature for Cd-doped Sn02. -473- 4 —3 min.-i pure - —< min. Sn0 a >—5 miOrJ A --3 min-i a , doped a—4 mi o —5 mMi Cd N~ 3 nn.J < 5t0 52 Fig. (7) -474- 3.2 Electrical properties 3.2.1 Sheet resistance Figure (8) represent the measured sheet resistance values of pure Sn02 films as a function of substrate temperature for different deposition times. A doped sample of substrate temperature 525°C and 10 minutes is also plotted in Fig. (8) for comparison with the pure samples. The Samples with 8% cadmium doping have lower sheet resistance values. In general, the curves show a gradual decrease in the sheet resistance values with increasing temperature and deposition times. The decrease in sheet resistance values in both pure and doped films may be understood as due to enhanced crystallinity which occures with increasing substrate temperature, as indicated by the x-ray diffraction data (Fig. 9 and Fig. 10). Beyond the optimum temperature (» 450'C) the sheet resistance show some increase. The possible reason for such behavior could be attributed to the neutralization of charge carriers by the diffusion of alkali material ions from the glass substrate. The same trend has been reported by other authors^23*2^. 3.2.2 Figure of merit * As mentioned before, all attempts are devoted to enhance the properties of the produced thin film either optically or electrically, Haacke et a!.t2°J put a ratio which 10 reflects these values, calling it figure of merit * = T /Rg, where T is the transmission percentage at the solar maximum (X = 500 nm) and Rg is the sheet resistance. Table 1 summarizes these values. -475 1E*7 A — 2 min -, * —3mia pure ~ 1E*6 o — A min. Sn02 o • — 5 min. *—lOmiaJ ~ 1E*5 "—10 mirv (doped a Cd6%) c 1E*4 t o in 'in 1E*3 i 1E*2 r 1E4.1 L 375 400 425 450 475 500 525 550 Substrate temperature ( C ) Fig.(8) Sheet resistance as a function of substrate temperature at different deposition time for pure and doped Sn02- •416. Table (1): The values of * (xlO'4) 3 min. 4 min. 5 min. Material US T% dooed Cd. 450*C 7.8 445 .90 12.4 350 .92 7.44 235 .84 500°C 4.8 670 .92 18 370 .96 9.61 200 .85 525*C 3.7 760 .88 10.7 450 .93 22 240 .94 It is clear from Table 1 that the highest values of figure of merit obtained for doped thin films. The large value of F allow these films to be used as an antireflection coating, as well as an electrode in solar cells. 4. X-ray diffraction X-ray diffraction patterns for pure and Cd-doped Sn02 films prepared at different substrate temperature in the range (400-525'C) are given in Figs. (9) and (10). It is clear from Fig. (9) that the films of Sn02 deposited at temperature lower than 450°C a broad hill and no peaks from crystalline material and observed which indicates the existence of an amorphous phase. But, the films of Cd-doped Sn02 (Fig. 10) prepared at substrate temperature 425°C shows crystalline structure and more than three intense peaks were observed. In general, the relative intensities of the diffraction peaks increases by increasing the substrate temperature. This indicates that the crystallinity increase withe increasing the substrate temperature for pure and Cd-doped Sn02. -477- (110) Ts = 525 'C c 3 (101) Ts * 450 C in c (211) (200) (301) rvw. Is = «5 C r~v i i_ 2A 30 i0 50 60 70 80 2 G (cleg) Fig. (it) X-ray diffraction pattern of pure Sn02 films prepared at different substrate temperature. -478- (110) (200) (211) Ts= 525 C (101) (220) (310X301) (101) (211) —~~>. (200) (110) Ts = 500 *C c (3 01) 3 O (101) (211) C i i, 21 30 40 50 60 70 8 0 2 6 (deg) Fig. (10) X-ray diffraction pattern of Cd 89c doped Sn02 films prepared at different substrate temperature. -479- The d-values calculated from the diffraction peaks are [25 in good agreement with ASTM data for Sn02 l However, the relative intensity of the diffraction peaks (I/IQ) does not agree well with the ASTM data. Agashe et al.'2"' reported similar observation in the case of pure Sn02 films. This is due to the fact that the films are textural and not composed of random grains as found in powder x-ray samples. The analysis of these diffraction pattern showed a polycratalline structure of tetragonal rutile structure^27!. The lattice parameters of these films are calculated and having the values : a = b = 0.4734 nm and c = 0.3182 nm, which are in agreement with the obtained by Manifacier and Fillard^27l General examination of the x-ray diffraction patterns for the doped films revealed that, the addition of Cd increass the intensity of the diffraction peaks. During the deposition some of the planes orient themselves to give maximum reflection, and hence maximum intensity is observed. Thus, the Cd-doped Sn02 films have better crystallinity. It has been also observed in the case of Cd-doping films at high temperature (525°C), two weak planes (220) and (321) were observed. This may be due to the effect of doping of Cd which increase the crystallinity of Sn02. An analysis was carried out on the preferred orientation of the planes in the investigated films by study the peak profiles of planes (110), (200), and (211), because these planes are more sensitive than in the (110) plane for describing the preferred orientation^2**.29!. It was observed that the diffraction peak intensities of (200) and (211) planes are more sensitive than in the (110) plane for describing the preferred orientation. Doping affects the preferred orientation of crystallines^^. From an analysis of the profiles peaks of Cd-doping films, it was found that the preferential orientation is along (200) plane. Afify et al.W reported that thin films of F-doped Sn02 were also preferentially oriented along (200) plane. -480- 5. Conclusion Spray pyrolysis technique has been used for the production of highly transparent, conductive, glass like, thin films of pure and doped Sn02 with different deposition time and substrate temperatures. Films with high optical transmission and low resistance have been obtained by doping 8% Cadmium. The substrate temperatures positiviely affects the conductivity. The pure and doped Sn02 films are polycrystalline and have the same structure. A change in preferred orientation is observed after doping. The large value of the figure of merit, *t>, in the case of doping film reflects the suitability of the film for different applications in solar cells. References 1. V.K. Jain and A. P. Kulshreshtha, Sol. Energy Mater., 4(1981) 151. 2. G.H. Heilmier, Sci. Am., 222 (1970) 100. 3. K.L. Chopra, S. Major and D.K. Pandya, Thin Solid. Films, 102 (1983) 1-46. 4. L.I. Maissel, in L.I. Maissel and R. Glang (eds.), Handbook of thin films technology, Mc Graw Hill, New York, 1970, Chapter 18. 5. J.P. Marton and D.A. Lepic, J. Electrochem. Soc, 121(1974) 2346. 6. J. Sam Maudes and T. Rodriguz, Thin Solid Films, 69 (1980) 183. 7. H.Pink, L. Treitinger and L. vite, Jap. J. Appl. Phys., 19(1980) 513. 8. W.A. bADAWY, H.H. Afify and E.M. ElGiar, J. Electrochem. Soc., 137(1990) 1592. 9. H.M. Afify, R.S. Momtaz,W.A. Badawy and S.A. Nasser, J. of Materials t. Scennce : Materials in Electronics, 2 (1991) 40. 10. J. P. Upadhyay, S.R. Vishwakarma and H.C. Prasad, Thin Solid Films, 167, (1988) L7J^10. -481- 11 E Shanthi, V. Dutta, A.Banerjee and K.L. Chopra, J. Appl. Phys.,• 1 (1980) 6243. 12 H Haitjema, J.J. Elich and C.J. Hoogendoorn, Solar Energy Mat., 18 (19890 283. 13. S Raghunath Reddy, A.K. Mallik and S.R. Jawalekar, Phys. Stat. Sol. (a) 96 (1986) K191. 14. A.K. Abbas and M.T.Mohammed, Phys. Stat. Sol. (a) 100 (1987)633. 15. m T. Mohammed and W.A. Abdul-Ghafor, Phys. State. SoUaO 106 (1988) 479. 16 C Agashe, M.G. Takwale, B.R. Marathe und v.G. Bhide, Solar Energy Materials 17(1988) 99. 17 RS. Momtaz, H.H. Afify and W.A. Badawy, In publishing. 18. EAktulga, Ph.D. Thesis, Dep. of Phys. Foe. of Sc. Istanbul Univ. Turky (1983). 19 H. Demiryont, I.R. sites and Geib, Appl. Opt, 24 (1985) 490. 20. G Haacke, Apple Phys. Lett., 28 (1976) 622. 21. O P. Agnihorri, B.K. Gupto and a.K. Sharma, J. Appl. Phys. 49(1978) 4540. 22 H Demiryont and K.E. Nittering, Sol. Energy Mater 1979,1989. 23. A.L. Unaoguw and C.E. okeke, Solar Energy Materials 20,29,1990 24. G.Sanon, R.Rup and A. Mansigh, Thin Solid Films 190, 28,1991. 25. Inorganic Index to powder Diffraction File, ASTM, PA, 50467. 26. C.Agashe, B.R. Marathe, M.G. Takawale and V.G. Bhide, Thin Solid Films, 164 (1988) 261. 27. J C. Manifacier and J.P. Fillarrd, Thin Solid Films, 77(1981) 67. 28. N S. Murty and S.R. Jawalekar, Thin Solid Film, 100(1983)219. 29 S.R. Vishwakarma, J.P. Upadhyay and H.V. Prasad, ibid., 176(1989)99. 30 A F Carrol andL.H. Slack, J.Electrochem. Soc, 123(1976) 1889, 1889. -482- First International Spring School & Symposium on Advances in Materials Science (SAMS 94) 15-20 March 1994 "Structural and Optical Properties of Cuo films prepared by spray pyrolysis" H.H. Afifi*, S.A. Nasser and S.A. El-Hakim Physics Department, Faculty of Science, Beni-Suef, Cairo Univeristy, Beni-Suef, Egypt * Physics Department, Solid State Division, National Research Centre, Dokki, Cairo, Egypt. Abstract: Cuo films have been deposited on glass substrate using spray pyrolysis technique. The films are single phase polycrystalline. X-ray diffraction peaks corresponding to the planes(110), (022). (Ill) and (111) increased in intensity with increasing time of deposition. Measurements show a good impovemem in the crystallinity with increas ing time of deposition. The optical parameter such as refractive index (n), the thick ness of the prepared thin film (t) and the absorption coefficient («) are calculatd from transmittancc and reflectance measurements between 1 ans 3Mm. The values of the direct allowed band gap depend on the time of deposition and varies between. 1.93- 1.7 cv. 1) Introduction : Copper oxides exist in two stable forms, the cuprous oxide Cu2o and Cupric oxide Cu 0. these two oxides have different colors, crystal structure and physical properties. Simple chemistry tells us that these differences arc mainly due to the fact thai Cu in + 2+ Cu20 is the Cu state, while Cu in Cu 0 is in the Cu state. Cu2o is a red-colored cubic semiconductor, Cuo on the other hand has a dark irongray color and crystallines in a more complicated monoclinic structure (1). The recent discovery of high temperature super conductivity in ceramic oxides (2) containing Cu has rekindled interest in studying the properties of copper oxides in much more detail. There arc many techniques for preparation of thin film Cuo (3-5). The spray meth od is particularly attractive because of its simplicity. It is fast, inexpensive and vac- cumlcss and suitable for mass production. This research aims to study the strctural, and optical properties of cuo prepare! sprey pyclysis in order to test how much this technique could lie used. 2) Experimental Procedure : The spray pyrolysis technique used is basically a chemical deposition technique in which the fine droplets of desired material solution arc sprayed onto a healed a healed substrate. A continuous film is pormed on the hot substrate by thermal decomposition -483- of the material droplets. The used spray cystem equipment is shown in Fig. (1) the de tails of the different compoment of (he system arc given in an prcvions work (13). Spraying 0.2 M solution of CuSo4 (Ricdcl - dc Hacn. W. Germany) in distilled water. The deposition parameters arc : air flow rate 5 ml min'1, substrate temperature 425 ± 2°C, and deposition time 10-25 min. X-ray diffraction was done with a phillips X-ray diffraction equipment model PW/ 1710 using Co K radiation. The optical measurements were made using the recording double-beam lambda 4 B Pcrkin - Elme Spectrophotometer. 3) Results and Discussion : 3.1. Structural Analysis: The films deposited at different times were analysed by X-ray diffraction. The val ues of a (interplancr distance) and the corresponding intensity of the diffracted lines 1/ lo (Table 1) confirmed the films to be made of crystalline CuO. It can be observed from X-ray diffraction pattern for the deposited CuO film at 425°C (see Fig. 2) of different times of deposition (10,15, 20.25 minutes) that the peaks corresponding lo (he planes (110). (002). (111) and (111) increased in intcrsily with increasing time of deposition. Therefore, a good improvement in the crystallinity is observed with increasing time of deposition. No preferred orientation is detected in X-ray diffraction patterns. 3.2. Optical Properties: The Alms deposited on chemically cleaned glass substrate arc examined using spcctropholomcntcr in the range 03-3 fim. Figs. 3 and 4 illustrate thr transmitlancc and reflectance spectra of the investigated samples. The general feature of (T-X) and (R-\) curves is the change of the percent transmitted and reflected light with time of deposition. The optical parameter refractive index (n) and the thickness ofthin Him (I) were calculated from the interference cxtcrma of the l-X iransmitancc spectra. The used method and equations for calculations the film thickness, absorption coefficient and refractive index has been reported elsewhere (6-10) The transmiltancc T is a function of absorption cocfricicnt («) and thickness (I) of the films as follows: 47rkt T = ACxp( - ) A WhereK=» X "- 4 7 2 2 2 2 and A = 16 na ng (N + K ) / <(na+n) + K ). ((ng+n)2+k2). -484- t being the thickness and n. na. n arc the refractive indices of the film, air and glass respectively. For K* «n\ A is approximately equal to unity and the exponential term domi nates, so that: Tr —exp(-«t) at the absorption edge. At lower energy, beyond the absorption edge, the absorption coefficient (<*) was calculated from (he expression « =L it, (-A. ) « Tr Where: 1 rr=-£i- (i. 2-» > (i + 22L )" C 2 Tmin Tmin. WilhCi=(n-n4) andC2 = (n-na)(ng-n) The refractive index (n) and thickness (t) were obtained front (7,8): , (n 2 + n 2) n' = ( • -2 ? +2nangt0) + (n ) 2 2 2 (( • ".*V +2nanet0) .na ne ))T andt= —- Tmax - Tinin where t„= Tmia x I nun N being the number of oscillations lictwccn two cxtcrma of the (Tr - A) curve, n,. and n2 represent the refractive indices at the two wavelengths A. i, and A. 2 respectively. The variation of refractive index (n) and the absorption coefficient (oe) arc shown in Figs. 5 and 6. To study the effect of the time of deposition of the prepared samples on the optical energy gap. the direct enccrgy gap of Cuo using equation (12): -485- « = o(hv-Eo)2 hv>Eo oc = o hv < Eo Where « is the absorption coefficient, hv is the photon energy and Eo is the ener gy band gap. The energy band gap is nearly constant and independent of photon ener gy- The absorption coefficient («) was carried out using the ratio of transmission or two films of diffeicnt thickness which is given by: rf T,.2 = c* Where Ti_2 is the ratio of transmission and t is the difference in thickness of the two films. « i can be determined without knowing t. This was achieved in the present investigation by placing the thin sample in the reference beam and the thicker ene in the sample beam of the recording spectrophotometer (13). The direct band to band energy gap is determined by plotting (« t)2 as a function of wavelength as shown in Fig. 7 Tabic (2) shows the energy gap values, it has been noticed that the optical energy gap of the prepared samples decrease with increasing time of deposition. References: 1- W.I. Ching. Yong-Nian X.U. and K.W. Wong. Phys. Rev. B 4011 (1989) 7684. 2- J.G. Bednorz. and K.A. Ma'er. Z. Phys. B. 64 (1986) 189. 3- C. Montcro. G. Poillcrat. Y. Rcas. J. Tcgada and P. Bosch. Sol. Energy Mater. 19 (1989)353. 4- M. Scrra and D. Sainz. Sol. Energy Mater. 13 (1986) 463. 5- A.B. Laurie and M.L. Norton. Mater. Res. Bull 24 (J989) 213. 6-J.C. Manifacicr.J. Gasiol andJ.P. Fillard. J. Phys. E. 9 (1976) 1002. 7- J.C. Manifiicicr. M. Dc Murica. J.P. Fillard and E. Vicario. Thin Solid. Films 41 (1977) 127. 8- S.I. Pankovc. Optical Processes in .semiconductors (prentice Hall. England Cliffs, NJ 1971). 9- O.S. Heavens. Optical propeiies of thin Solid Films (Dover. New York. 1965). 10- T.S. Moss. Optical properties of Semiconductors (Bulicrworths. London 1959). -486- 11- A. Mondal. S. Chaudhui and A.K. Pal. Appl. Phys: A. 43 (1987) 81. 12- A.K. Abass, H. Baker, S.A. Jassim and I.A. Fahad. Sol. Energy. 13- H.H. Anfy, R.S. Momiaz, W.A. Badawy and S.A. Nasser. J. Malcr. Science. Mal cr in Electron. 2 (1991) 40. -487- Table I: Comparison of the observed il-values of ile|Kisited CtiO films wild ASTM curd rc|X)rtcd. Time of Temperature (J(Ao) d(Ao) Inlcrsity Intensity deposition of prepara observed ASTM of Miller Index observed reported (minulcs) tion oc. CuO powder F/io I/Io 10 425 2.775 2.751 110 30 12 2.559 2.530 002 36 49 2.512 2.523 III 100 100 2.316 2.323 111 100 % 15 425 2.783 . 101 30 . 2.754 2.751 110 20 12 2.559 2.530 002 30 49 2.518 2.523 III 100 100 2.516 2.323 III 92 96 20 425 2.790 . 101 40 . 2.734 2.751 no 16 12 2.559 2.530 002 28 49 2.512 2.523 III 85 100 2.326 2.323 III 100 % 25 425 2.776 . 101 25 . 2.740 2.751 110 14 12 2.565 2.530 002 30 49 2.518 2.523 in 100 100 2.316 2.323 111 98 96 Table (2) Optical Energy Gap of CuO Time (min.) 10 15 20 25 Energy 1.93 1.82 1.87 1.70 gap (ev) -489- Fi-j. (2) x-ray diffraction patterns of Cu o films prepared dt different tiroc* of deposition a) 10 *iin ti) IS nan c) 20 nln d) 'i'j »m -490- uoo jcoo cm KKJ. (J) Vinn±ndti:tnc.ti Tt ct' Coo ti.ju fii.-.i of cliff«r<.r.« tii.ic of iii?[josition u) 10 Min b) 15 Mir» e) 20 Min a) 25 Man m> I.**** / .VI..IIUIU a^l«IU wi ll*% VI4f |I»I*4]«M •|«WIK -491- -491- to 2 So U2i 3000 crtr Fig. (-1) Kefiectance Kfc of CMO thir. films of different time o>. deposition «*) 10 Min b) 15 Min c) iO Min d) 25 Min V (a) li \ /C\ a) X («" ^^ co s us 1 , 1 i *-s •25 Fig. (5) Refractive index , as a function of wave lenyth fur CM O thin films of different time of depo sition a) io Min b) 15 Min c) 20 Hin d) 25 Hin o Fig. (6) The dependence of LoglQ of the absorption coeffeclent on the wave length for Cu o thin filrs of different time of deposition a) 10 Kin b) 15 y.in c) 20 Kin a) 25 Kin 4* Fig. (7) Variation of the equaii of the absorption coefficient as a funtlon of wavelength of Cuo thin films of different time of deposition a) 10 Min ' b) 15 Min c) 20 Min d) 25 Min -495- First International Spring School & Symposium on Advances in Materials Science (SAMS 94) 15-20 March 1994 Effect of Modified Evaporation. Source on The Structural Characteristics of CdS Films A.Ashour, N.El-Kadry and S.A.Mahmoud Physics department, Faculty of Science, Mini* University, Minia, Egypt. In this work CdS films were formed by physical vacuum deposition technique onto an amorphous (glass) substrate using modified evaporation source. Effect of processing conditions (substrate temperature, deposition rate and film thickness) were considered in comparison with the normal evaporation source. The structural features and morphology were investigated by scanning electron microscopy, and the microstructures by X-ray intensities and line profile analysis (Warren-Averbach method). Microcracks were observed due to the difference of thermal expansion between the film and the substrate. No piles were observed in contrast to the case of normal source. Using a shutter to cover the substrate during the evaporation starting-up and covering the source with quartz wool is needed to improve the surface defects of the films. The X-ray diffractograms show a highly crystallographic oriented films with the hexagonal c- axis perpendicular to the substrate. The crystal perfection of the films are improved by lowering the deposition rate or by ascending the substrate temperature and the film thickness. 1. Introduction. CdS thin films have been employed by various workers as semiconductors in the insulated-gate-thin-films field effect transistor, known as the TFT[1,2], in heterojunction[3,4] and optoacoustic transducers[5]. These devices are most commonly made by vacuum deposition of the materials onto an amorphous substrate such as glass. A knowledge of the microstructure and electrical properties of semiconductor layers used in the thin film solar cell is essential in order to form a realistic model for the device and to improve its efficiency performance, reproducibility, and reliability. In the present work, the structural properties of vacuum deposited CdS films formed under various preparation conditions will be investigated. The structural characteristics include phases present, crystallite size, residual microstrain and nature of preferred orientation. The measured structural characteristics will be intercorreiated. 2. Experimental The evaporation of CdS (5N purity, Fluka AG) was earned out in a conventional vacuum system (Coating unit E306A Edwards Co.) prowded with an -497- oil diffusion pump, via a liquid nitrogen trap. The lowest attainable pressure was in the range (1.3-2.0)X10"3 Pa. The source was a cylindrical necked quartz crucible wrapped in a tantalum sheet. This arrangement was found to facilitate the loading of the CdS powder and to increase the lifetime of the heater. A small wad of quartz fiber wool was placed in the neck of the quartz crucible to prevent spattering of the Cdo powder during outgassing and evaporation. The substrate holder was made of copper sheet-provided with a heater and the temperature was measured by copper-constantan thermocouple and was controlled during deposition. The film thickness and the deposition rate were monitored using standard quartz crystal techniques. Structural analysis of the films was carried out by scanning electron microscopy using a Jeol Stereoscan electron microscope (model ST200). A Jeol X-ray diffractometer (model 60PA), using filtered Cu-ka_radiation was employed to obtain the diffraction patterns of the films deposited on glass substrates. For this study, a slow scanning speed and a small time constant were applied in order to attain intensity of high accuracy. 3. Results and discussion Scanning electron micrographs are shown in Fig.[1]. Cracks are observed which are probably present because of the elastic strain produced in the film after deposition, due to difference in thermal contraction between the deposited film and the glass substrate. Similar cracks were observed by Ashour et al.[6] for CdS films prepared by the other type of evaporation source to produce CdS/CdTe solar cells. It is worth mentioning that the piles of CdS materials which have escaped from the boat during evaporation and adhered to the surface of the film[6,7] were completely absent when the present modified source was used. Figs.(2-4] show the diffractograms of the prepared CdS films and that of the randomly packed powder is given in Fig.[5] for comparison. Fig.[5] is a typical X-ray pattern of the wurtzite structure of CdS. However, the films show a highly oriented crystallite with the hexagonal c-axis perpendicular to the substrate. The analysis has shown that the diffraction patterns of the films are essentially of the (OOI) planes. This indicates that the films grow, under all the preparation condition studied with the * jt ^ J-1 tray;' • • -h '}•><• v • -499- (COT) (*aa> - (CBT) * I "I a a • a (Zaoi ~ OKDW1M \ *sts tatfts l»T«f» 'OJW)A»T«U«»UI ilia Uiiii OJHHHri (est) ugiiiii as • • • • (sat) - a ffb. »!i (zoai •V r. «C OTUfl 'q^W)AaT»u»3UX -500- (can (Ctttl « S BESS esc! aNCN X ONOM 44 craa» Si " s«ssss ^V- 5 5:552555! a s Sliliiiim* U 11 H It *j a A.» w ** *J fiSiJLJl S5=2S£, «wti 9 lb i> J ICBtl s •9 « 8 s * * s fat»n •«jv>«^tsua>«i -50!- I.?0 a.ca CM *•'* It) •.<• ». 1» 1' t.it c a t.it l.tt 0 t. te 13} t a . . _ t.JI > > <: 1 s ... c Id il c 1. «• A j,» 1 i.jt * i.it c( *f I.Ct JE l.tt 1.68 t.M 141 t.M (31 ft 9. • • f i -» 1 .«• o O n n a s a s B.20 1 8 i •4 O 5 1 . 1 0 1 u.» «• II.» Jt.t t.« j«.o w.e it t H i ZOdJvq.l ZBIUoy.l Fig.141: Effecc of film chickness on X-ray diffracccgrams of CdS films. •- (11- t=10Q0 nm. Tg=l80 °C. R=O.B iWs (4>- t = 1500 run. Ts=180 °C. R=0.8 nin/s 13)- t»lOOO nm. Ts=180 «C. R=2.2 nm/s (S»- t=lSG0 nm, Ts=180 °C. Rs2.2 nm/s - •OT lit * • < C.83 ...... I 1 ' < a, •OT 8 •% • H» 0 f» •OT 1 r- - Mm, urn « - -503- Table[ 1]: The values of the crystallite size and the microstrain of CdS films. , jSample T$ t R GS Strain | jNo. (°C) (nm) (nm/s) (nm) (10'3)| | (1) 180 1000 0.8 80 1.01U J ' (2) 220 1000 0.8 164 0.658 j .I (3) 180 1000 2.2 39 1.568 j [ (4) 180 1500 0.8 147 1.032 ! ; (5) 180 1500 2.2 52 1.297 J | (6) 200 1500 2.2 82 0.813 J The crystallite size and the residual microstrain have been evaluated by the Warren and Averbach method[8]. The crystallite dimension was found to increase with the increase of film thickness and substrate temperature but to decrease with increasing deposition rate. The values of the crystallite size and the microstrain of CdS films are given in table[1]. As can be seen from this table, the crystallite size values are larger than those obtained when the normal evaporation technique(molybdenum boat) was used[6]. On the other hand, the microstrain values are smaller. In fact the larger values of the grain size corresponding to the increasing of the stoichiometry and the film continuity when used the modified method. The decrease in the strain values is attributed to the increase of the grain size and the absence of piles. The crystallite size increases as the substrate temperature and the film thickness increase or the deposition rate decreases. However, the behaviour of microstrain is vice versa. Such behaviours are in agreement with previously published data[6,9]. films deposited at high Tg have a low density of initial nucleation sites and the resultant is large crystallite size. In addition, there is a tendency for a grain growth and lattice defects annihilation as T5 increases, which also results in the observed increase of the crystallite size and relief of the internal strain. The increase of the crystallite size is the normal behaviour observed for materials in thin film forms, whatever is the evaporation source. One of the discrepancies in the literature is the effect of deposition rate on the internal strain in the film. Usually, the contamination and the internal energy play an opposite roll. In this work, we believe that high deposition rate results in the increase of residual strain. Conclusion The structural investigation by SEM and X-ray analysis prove that the used modified evaporation source (heated quartz tube covered with a small wad of quartz fiber wool) produces film of large crystallite size and lower internal strain than the normal evaporation source (Molybdenum boat). In general, the crystallite size increases as the substrate temperature and the film thickness -51)4- increase. On the other hand, the crystallite size increases with decreasing the deposition rate which relieves the internal strain within the film. Acknowledgment The authors would like to thank Prof.Dr.A.A.Ramadan, Physics Department, Minia University, for his help and useful discussions. 5. REFERENCES [1] P.K.Weimer, Proc.l.R.E, 50, 1462, 1962. [2] R.Zuleeg, Solid-Slate Electron., 6,193,1963. [3J A.L.Fahrenbruch, J.Crystal Growth, 39, 73,1977. [4J K.Yamaguchi, N.Nakayama, H.Matsumoto. and S.lkegami, JapJ.Appl.Phys., 16,1203, 1977. [5] S.S.ELLiot, V.Domarkas, and G.Wade, IEEE Trans.Sonics Ultrasonics, 25, 346,1978. [6] A.Ashour, R.D.Gould and A.A.Ramadan, Phys.Stat.Sol.(a), 125,541,1991. [7] A.G.Stanly, Appl.Sol.Stat.ScL, 5, 251,1975. [8] H.P.Klug and LE.AIexander,"X-ray diffraction Procedures for Polycrystailine and Amorphous Materials", 2nd ed., Wiley, New York 1974. [9] H.H.Afifi, S.A.Mahmoud and A.Ashour, The XVI conference on solid state Science, Minia 9-14 Jan., 1993 -505- EG fScMS 23 First International Spring School &. Symposium on Aduances in Materials Science (SAMS 94) 15-20 March 1994 Advanced Dilatometrie Technique "SID" for Studying the Kinetics of the Intermediate Stage of Sintering. Example: Sintering of ITO, and CeO, M.El-Sayed Ali. O.T.Sdrensen' & S.El-Houte Metallurgy Dept., NRC, Atomic Energy Authority Cairo-Egypt * Risyi National Lab., Denmark. ABSTRACT The stepwise isothermal dilatometry (SID) Is extended to study the shrinkage kinetics at the intermediate stage of sintering. Based on Coble's model and the linear dependence of the cube of the grain diameter on time; an equation describing the intermediate stage of sintering, has been derived as function of shrinkage and shrinkage rate. The intermediate stage shrinkage kinetics of UO, compacts - sintered in H, and H, * 10% CO, and CeO, compacts sintered in oxygen, have been studied. Reasonable values of activation energies of 516.2 ± 4S. 438.9 ± 35, 422 * 34, KJ/mole have been obtained, respectively. These satisfactory results reveals the validity of this technique in studying the intermediate stage of sintering. I-INTRODUCTION The intermediate stage of sintering is-in fact- a very important stage of the sintering process, because it extends from the end of the initial stage (at about 5* shrinkage) up to 95% of the theoretical density (1). Contrary to the initial stage where the grain growth is almost absent and has no significant effect on the shrinkage; the intermediate stage is characterized by simultaneous grain growth and densification. Coble (2) introduced the first model describing the intermediate stage of sintering, based on the Kelvin's tetrakaidecahedron a tedious procedure for calculating the mean average grain intercept length, using sample impregnation, sectioning, polishing, etching and optical examination. Semiempirical (4) and statistical models (5) , have been also proposed employing a simplified geometry (6). Due to the interplay of various parameters, simplifying assumptions have to be made. However, none of them has been fully accepted. On the otherhand Coble's model (2), based on sound theoretical considerations, regarding -5(17- the grain geometry, showed excellent agreement with experimental results obtained for the intermediate stage of sintering (7). Attempts have been made to study the intermediate stage of sintering, using a constant heating rate dilatometric technique (8,9), the results have been analysed according to Coble's model: The stepwise isothermal dilatometry (SID) is a relatively new technique, by which the initial stage of sintering has been studied successfully (10-12). This was achieved by performing the measurements - using only one sample in a single run. The (SIDi technique has been used recently in a practical application, by which the powder compact's sinterability could be evaluated (13). Lately a computer controlled (SID) version has been developed for a faster shrinkage kinetic determination (14) . In this work, the (SID) technique will be used to study the intermediate stage of sintering of U02 and Ce02 powder compacts. Coble's model will be used to derive the kinetic equation, by which the shrinkage data will be analysed. Principle of the stepwise isothermal dilatometrv (SID) The principle of the (SID) have been described in a previous work by the authors (14). In brief, it is based upon using the shrinkage rate to control the sample heating, so that isothermal segments could be obtained at different temperatures. This is illustrated schematically in Fig.l. The advantages of the (SID) is obvious, i.e. instead of performing three or four dilatometric investigations with different samples, as shown by the dashed lines on the Fig.l, the (SID) can create isothermal segments at the corresponding temperatures -necessary for the kinetic analysis- in a single run, using only one sample. The mathematical formulation for the (SID) during the intermediate stage of sintering: For shrinkage data analysis, in this section an equation, that can be used for shrinkage data analysis, during the intermediate stage will be derived. This might be obtained on the basis of Coble's model for cylindrical pore shrinkage, by bulk diffusion as expressed by equation [1], dP > K D.,JL£L = K(T) 1 [13 ST d1 RT ~S? where K = constant, Dv = volume diffusion coefficient, p - surface energy, d » average grain diameter, R = universal gas constant, T - absolute temperature, K(T) = temperature dependent constant. By substituting for the grain size using the following equation (15) . d> = dj + K' (T) .t where d is the grain diameter at t=0, K' (T) = a temperature dependent constant. then, integrating equation [1] we get : -508- P =P„ - A. In t d' • K' (T) .t] [2] where A - K(T) = cosntant,assuming equal temperature dependence, KMT) i.e the activation energy for grain growth is consistent with that for densification (16). Equation (2] is then differentiated to give: P = dP « - A. K' (T) exp - (P„-P) [3] dt A Porosity P could be easily expressed by a function of the shrinkage as follows, where P , P,h are the initial and theoretical relative densities, and Y is the relative shrinkage. putting S= P / p and X» 1 then. X= 3Y and P--SX 15) (1-Y)' substituting from [4] and [51 in [3] one can get In X = In K"(T)-S/A (X-l) [6] where K"(T) • A./S K'(T) is the temperature dependent constant for the shrinkage kinetic of the intermediate stage of sintering that are characterized by the grain growth .and shrinkage. Plotting data according to equation [6], { In X vs. (X-l)],a set of parallel straight lines are to be obtained. Each line represents an isothermal segment. The intercepts with In X axis, gives values of In K"(T), which if plotted versus 1 / T gives a straight line, from whose slope, the activation energy could be calculated. II-EXPERIMENTAL TECHNIQUE Nuclear grade uranium dioxide and 99.9 * purity Ce02 powders, were used in this study. The powders had BET specific surface areas of 5.97 and 4.59 m2/g respectively. Cylindrical compacts were prepared by pressing the powder at 100 MPa. The green densities were 54 and 61V of the theoretical densities respectively. The measurements were performed in a horizontal dilatometer (Netzch Type 402 E) with alumina sample holder. The sample heating was computer controlled, taking in consideration, the shrinkage rate and its preset limits (14). which are predetermined in order to obtain enough number of isothermal segments (reasonably long in time scale), that are necessary for performing the kinetic analysis at each sintering stage. The linear variable differential transformer (LVDT), gives the length change signals via an analog digital converter to the computer, which in turn is used to control the sample heating •5W- and -of course- to be the source of the shrinkage data at the same time. The UO, compacts were fired in hydrogen atmosphere during the first experiment, while in the second, hydrogen plus 10V C02 atmosphere was used. The CeO: compacts were fired in oxygen atmosphere. Ill-RESULTS AND DISCUSSIONS The shrinkage-time data were polynomially fitted and the corresponding shrinkage rates were obtained. Then the parameters X and X were easily calculated by the computer. Plotting In X vs. (X-l) according to equation [6]. a set of nearly parallel straight lines were obtained at different temperatures. Fig.2 shows the results obtained upon sintering of U02 in hydrogen atmosphere. Similar results are shown in Fig.4 and 6 for the sintering of U02 compacts in hydrogen containing 10% carbon dioxide,and Ce02 compacts in oxygen, respectively. The above results show clearly to what extent the experimental results fit the derived equation. Consequently, the results suggest that Coble's model for the shrinkage of cylindrical pores by bulk diffusion, can express quite well the shrinkage kinetics in the intermediate stage of sintering. The slopes of the parallel lines are comparable, in case of sintering U02 compacts in H2 + 10% C02 because of their equal green densities; which means the same value for S/A. As for Ce02, higher values are obtained due to its relatively higher green density compared to that of U02. rne temperature dependent parameter In k" could be obtained from the intercepts of these parallel lines with the In X axis. Arrhenius plots are obtained by plotting In k" vs. 1/T. Fig.3,5 and 7 show the straight lines of these Arrehenius plots. From their slopes, activation energies of 516.2 ± 46, 438.9 ± 35 and 422.2 ± 34 KJ/mole were obtained respectively. The slope and the correlation coefficients were obtained from linear regression analysis. The activation energy obtained for the intermediate stage sintering of U02 compacts in hydrogen (516.2 KJ/mole), is relatively higher than the one obtained for the initial stage sintering; 390.0 KJ/mole (11). This might be an indication of the existence of different operating mechanisms. On the otherhand, the first value is in good agreement with the values obtained from the grain growth kinetics of U02 compacts -518 KJ/mole-, in which the cubic dependence of the grain diameter as a function of time was best represented (15). The activation energy obtained upon sintering of 00, in H2 + 10% CO, atmosphere is much higher than that obtained from the initial stage: 234 KJ/mole (10). This can be attributed to chanqes in the equilibrium towards more reducing conditions, which results in reduction in cation vacancies; this occurs at higher temperatures (17), As for the case of Ce02. - slight decrease in the activation energy, relative to that of the initial stage (12), was observed. -510- CONCLUSIONS 1- The (SID) technique could be used to study the intermediate stage sintering as well as the initial stage, in one single experiment, using only one sample. 2- Based on Coble's model and the cubic dependence of the grain diameter on time, a mathematical equation has been derived, which successfully represents the experimental results. 3- Activation energies of 516.2 ± 46, 438.9 ± 35 and 422,2 ± 34 KJ/mole were obtained for the intermediate stage of sintering of U03 compacts in H,, H2 + 10%CO2 and ceOj in oxygen, respectively. REFERENCES 1- R.L. Coble, "Sintering Crystalline Solids. 11. Experimental Test of Diffusion Models in Powder Compacts," Journal of Applied Physics, 32,5,793-99, 1961. 2- R.L.Coble,"Sintering Crystalline Solids.I.Intermediate and Final State Diffusion Models," Journal of Applied Physics, 32,5,787-92, 1961. 3- D.L.Johnson."A General Model for the Intermediate Stage of Sintering," J.Am.Ceram.Soc., 53,10,574-77, 1970. 4- J.H.Rosolowski and C.Greskovich,"Theory of dependence of Densification on Grain Growth During Intermediate-stage sintering,"J.Am.Ceram.Soc.,58,5-G,177-82, 1975. 5- T.Ikegami and Y.Moiyoshi,"Intermediate-Stage sintering of a Homogeneously Packed Compact," J.Am.Ceram.Soc.,67,3,174-78, 1984. 6- D.L.Johnson,"Solid State Sintering Models," The Sintering Processes, Material Science Research, Ed. by G.C.Kuczynski, vol.13,97-106, 1980, Plenum Press. 7- T.K.Gupta,"Comments on Sintering Kinetics Based on Geometric Models," J.Am.Ceram.Soc., 52,3,166-67, 1969. 8- R.M.German,"Grain Growth Influences on the Sintering Densificstion of FCC Metals, Example of Palladium," Sintering Processes, Science Research, Ed. by G.C.Kuczynski, vol.13 Plenum, 1980. 9- C.Genuist and F.J.M.Haussonne," Sintering of BaTi03: Dilatometric Analysis of Diffusion Models and Microstructure Control", Ceramics International, 14.169-79, 1988. 10- M.El-Sayed Ali, O.T.Sorensen and L.Halldahl," Quashi- Isothermal Dilatometric Studies of the Influence of Oxygen Pressure on the Initial Sintering Stage of U02 Powder Compacts, " Journal of Thermal analysis, 25, 175-80, 1982. 11- M.El-Sayed Ali, O.T.Sorensen and L.Halldahl, "The Vse of Quasi Isothermal Dilatometry in Evaluation of the Initial Stage of Sintering U02 Powder Compacts", Thermal Analysis: Proceedings of the Seventh International Conference on Thermal Analysis. Ed. by, B.Miller, Wiley, 1982. 12- M.El-Sayed Ali and O.T.Sorensen. "Initial Sintering Stage Kinetics of Ce02 Studied by Stepwise Isothermal Dilatometry", Riso-R-518. 1985. 13- M.El-Sayed Ali and O.T.Sorensen,"Practical Application of Stepwise Isothermal Dilatometry for Characterization of Sinterability of Powder Compacts", Riso-M-2694. 1988. -511- 14- M.El-Sayed Ali, S.El-Houte and O.T.Sorensen, "Computer Controlled Stepwise Isothermal Dilatometry", to be published in Interceram. 15- F.A.Nichols, "Theory of Grain Growth in Porous Compacts", J.Applied Physics, 37,13,4599-602, 1966. 16- R.L.Coble and J.E.Burke, "Sintering in Ceramics", Progress in Ceramic Science, vol.Ill, Pergamon, London, 197-250, Ed. by, J.E.Burke, 1963. 17- M.El-Sayed Ali and R.Lorenzelli, "Kinetics of Initial Stage of Sintering of U02 with Nd203 Addition", J.of Nuclear Materials, 87, 90-96, 1979. -512- IT v> Pig.l Principal of (sin) LN X -1U*- -»U f- -n« •nj Fie.2 In x v«. (x-1). obtained for the shrinkage of UO, compacts sintered in H,. at different teaperaturea. B JO«l*C • 10»J*C * IIU'C w i:s-7*C • »9«#C a J354*C in K(T) M .88 .73 «00/T(K*) Flg.l Arrehenlua plot for K-(T) vs. l/T for sintering of UO, compacts in H,. -514- -TOM -«U -IU - •tu ri«.4 la i »«. U-*l. obtained tor eh« shrinkage of 00, caapaece, sintered in N, » 10% CO.. at different teaperatiureta: U>K *0B/T(»f) Fie.S Arrehenlus ploc for It-(T) vs. l/T for sintering of UO, coapactsjla K • 10% CO,. -515- •TUMLX •tut \- -12.4 Fly.? In x vs (X-l). obtained for the shrinkags of ceo, compacts, sintered in O,. at different temperatures: . -4 i -»- c««;/oa. o-taa KJ/IDOK coo. co»t o.*»» -e r -7 -W .M .87 .71 .78 «00/T(k) Fig.7 Arreheniua plot for K'.,-H vs. l/T for sintering of 00, compact in O,. -516- First International Spring School & Symposium on Advances in Materials Science (SMfS 94) 15-20 March 1994 THE USE OF ALUMINA OR MULLITE IN FABRICATING SILICON CARBIDE COMPOSITE MATERIALS M .A. A. Elmasry *, and A.M. Gadalla **. ABSTRACT The effect of additions of AI1O3 powder or AI2O3 produced by AI(N03)3 . 9H;0 hydrolysis to fine and coarse SiC powder on the green density of CIP bodies was investigated . It was found that the green density increases by addition of AI1O3 to SiC and decreases by precipitating Al(OH)3 on SiC due to the coaling of the latter by AHO3 after calcination. Pressureless sintering of SiC- AI2O3 composites is accompanied by a. fight loss which increases by increasing the firing temperature or the mole% alumina up to 30% and decreases above this ratio . The weight loss was attributed to the formation of the gaseous products SiO, CO, and AhO which are volatile at the sintering temperature used . Aluminum oxycarbides and aluminum carbide were identified in the sintered compacts. while in some samples SiC disappears or residual traces are found. Mixtures containing SiC and mullite showed a weight loss after pressureless sintering higher than the corresponding mixtures containing SiC and alumina under the same conditions. Hipping retarded the weight loss due to the inhibition of the reactions causing the weight loss. It was found that hipping was more effective in reducing weight loss in the case of SiC - mullite composites than in the case of SiC - alumina composites. The compatible phases were used to construct the tie tetrahedra in the Si - C - Al - O system and hence in obtaining the AI1O3 - AI4C3 - SiC - Si02 quaternary diagram. * Metallurgy Dept.. NRC, Atomic Energy Authority, CAIRO, EGYPT . ** Chem . Ens • Oepl., Texas A & M Univ., College Station , Tx.. USA. -517- I. INTRODUCTION: Development of structural ceramics for use in heat engines allows operation at higher temperature . reduces imported strategic materials and saves energy. Friction is expected also to decrease . and if 50% reduction can be achieved .elimination of lubricating oil will be possible. Moreover , ceramics will yield a reduction in the engine weight and elimination of cooling system in piston engine. Silicon carbide is a candidate material for the high temperature parts and is used in producing some composites (I), and in preparing prototypes of diesel and advanced gas turbine engines due to its high thermal conductivity, high elastic modulus, low thermal expansion and accordingly high thermal shock resistance (2,3). New generation of materials with high toughness which can be operated at higher temperatures in severe environments are required. The DOE and NASA support car manufacturers . ceramic producers and several research laboratories in carrying out research using SiC and Si3N4 whiskers and fibers in different matrices. Work by Bescher's group is an example to show the development in this area (4). Alumina matrix attracted the present authors since the diffusion of oxygen in alumina is extremely low (5) . and is expected to protect SiC from active and passive oxidation in severe conditions. Muilite attracted our attention since it was found to adhere to SiC and S13N4 bodies (6,7) . and since it may form as an oxidation product (S.9). Coating SiC with muilite prior sintering in alumina matrix may be beneficial because it has intermediate mechanical properties between SiC and alumina. Since most of the material prepartion , in published work , was done under high pressures and since the products will be used under low as well as high pressures in low oxygen potentials , this study concentrates on characterizing reactions occurring under normal pressure in low oxygen partial pressures, methods to retard them and their effect on densiflcation . tt is to be noted that Jackson et. al. (7), reported that some volatilization took place during the sintering of a or p - SiC at 1850 °C - I950°C in Ar, and they attributed the weight loss to vaporization of the liquid phase between AbC»3 , ( which they used as a sintering aid ) , and also to the reactions between AI2O3 and SiC , but they did not give any values or figures for the weight loss, although they recommended short sintering periods to limit this weight loss. The only reaction they gave , which they believe it was responsible for SiC decomposition and hence the weight loss was: SiC + Al203 = SiO(g) + Al20(g) + COfg) -5IX- 2. EXPERIMENTAL PROCEDURE : The a - SiC used was supplied by Norton . contains 0.15 . 0.63 . 0.36 . 0.08 . 0.08 . 0.05 . and 0.03 % weight Si - free . silica . C - free . Fe . Al. CaO. and MgO respectively. The purity was 98.65 % weight. The average particle sizes were 23 um and 5.5nm.Thect - alumina was supplied by Johnson Matthey. The purity was 99.99 % weight and the average particle size was 1.5 urn . The aluminum nitrate was supplied by Fisher with AKNC^^H^O,101.3 % weight and 0.0036 impurity as stated by the supplier. The alumina powder was mixed with coarse and fine SiC powder ultrasonically in isopropanol followed by drying at 80 °C . In another series of experiments finer alumina was precipitated from a solution of AKNC^^HoO using NH4OH at pH between 9 to 10. The slurry was filtered , dried at 200 °C and calcined at 400 °C for lh . SiC - mullite composites were prepared using two techniques . In the first technique the mullite was prepared from AI(NC»3)3.9H20 and SKOCiH^ by fast hydrolysis using NH4OH to adjust the pH (4,5) This was followed by drying at 60 °C and calcination at 1200 °C for lh . The mullite prepared by this technique was wet mixed with SiC ultrasonically . In the second technique the mullite was prepared from Al(OC2H5)3 and Si(OC2H5>4 using HMO3 as a catalyst and water for hydrolysis while agitating with SiC powder (6). The thermal analysis for the mullite showed a peak at 966 °C for the firsttechniqu e . while that for the second technique was at 970 °C . This indicates the good mixing of the raw materials and also the formation of mullite (5). The powders prepared by the above techniques were cold isostatically pressed (CIP) at 55K PSI to prepare the green bodies , which were then placed in a graphite crucible , surrounded by SiC powder and covered with a graphite lid . Firing was then carried out up to 2125 °C in a graphite furnace using stagnant helium atmosphere . The effect of pressure on reactions was followed by firing at 1825 °C in a hot isostatic press up to pressures of 28K PSI. Green and fired densities were determined by measuring the bouyancy when immersed in mercury . Fired discs were characterized microscopically and by X-ray diffraction to identify the phases present. The oxygen contents of argon and helium were zero % . -519- 3 . RESULTS AND DISCUSSION : The green and fired relative densities for SiC - AI2O3 composites from 10 to 85 mole % AI2O3 are shown in Fig. (1). These densities are represented relative to the theoritical densities calculated for each mixture The variation of green densities obtained by wet mixing of 23 urn or 5.5 urn SiC with 1.5 urn AI2O3 was as expected from mixing coarse with fine particles . While addition of fine panicles to coarse ones causes the density to increase due to filling the interstitial voids between coarse ones . addition of coarse particles to fines causes the density to increase due to the substitution of fines and their voids The higher the ratio between coarse and fine particles the higher is the density of the mixture . Accordingly for the same chemical composition , coarse SiC produced higher density , Fig.( 1) On the other hand, formation of AI2O3 from A1(N03)3.9H20 behave differently and the green density decreases as more fine particles precipitate. Moreover, the densities obtained by precipitation were much less than those obtained by wet mixing. This effect was attributed to coating SiC instead of filling the voids. To confirm this conclusion elemental distribution maps , after calcination , were obtained and all SiC particles were found to be coated completely with alumina . Only very few small particles were found to contain no Si . In such particles SiC may be completely masked by the precipitated alumina. The percent relative densities obtained for the cold isostatically pressed composites fired at 1825 °C for 4 h are shown in Fig.( I), which indicates that increasing the AI2O3 leads to an increase in the fired density. This result can be attributed to the fact that the sinterability of AI2O3 , which has high ionic character, is much higher than that of SiC which has a covalent character, and has a low surface to grain boundary energy ratio (1). The low green densities obtained due to coating are responsible for the low firing densities in mixtures prepared by alumina precipitation . The authors initially thought precipitation yields active and higher sinterable materials . In some specimens a remarkable toss in weight was obtained and at slightly higher temperatures some specimens were highly porous and others nearly vanished . Accordingly the weight loss of the composites was followed up . At the same temperature , Fig.(2), the weight loss was found to increase by increasing the amount of alumina up to 30 % , above which weight loss decreases . Loose SiC and AJ2O3 powders were heated under the same conditions in nearly pure argon (99.998 % purity), and gave only 3.3 and 0.0001 % weight loss respectively . The loss in weight may be due to the decomposition of SiC and the formation of Al<>0, AIO , or AlCb in the case of alumina . To study the relative stability of aluminum and silicon oxides , the standard free energy changes for the reactions containing these oxides were calculated and plotted against temperature . The results are shown in Figures (3) and (4). while AG for the reaction : 2Si02 + Al20(g) = 2SiO(g) + Al203 (I) was found to be negative above 500°K . those for the following reactions : SiCb + 2AIO(g) = SiO(g) + Al203 (2) -520- Table (1) Phases identified by X - ray analyses for SiC - Ab O3 and SiC - mullite composites. Sample Composition . C mole%) Temp., Time, Phases identified by X-ray No. SiC Ah03 mullite Si02 °C li 1 15 85 . . 1825 4 a - AI1O3 . SiC .AI4O4C 2 30 70 - - 1825 4 a - AI1O3 SiC,AI404C . 3 50 50 - - 1700 1 a - A1203 , SiC,AI404C 4 45 55 ~ ~ 1800 8 a - AI2O3 , SiC, AI2OC + (traces of AI404C). 5 45 55 " " 1825 4 « -A12O3.S1C.AI2OC + (traces of AI404C) - 6 60 40 * 1800 12 a - Al203 , SiC ,AI2OC + (traces of AI404C) 7 60 40 - - 1825 4 a - Al203 , SiC ,AI2OC 8 75 25 - - 1800 12 a - AI2O3 , SiC ,AI2OC 9 75 25 - - 1900 4 a - Al203 , SiC ,AI2OC 10 75 25 - - 1825 4 a - Al203 , SiC ,AI2OC 11 80 20 - - 1825 1 SiC, AI2OC. 12 90 10 - - 1900 4 SiC,AI2OC. 13 90 10 - - 1825 4 SiC,AI2OC. 14 7.8 • 92.2 - 1675 1 a -AI2O3 . AI2OC . mullite. 15 21.1 - 78.3 - 1675 1 a -A1203 , AI2OC . 16 50 - 50 - 1775 I a -AI2O3 , SiC , mullitc . 17 60 - 40 - 1825 0.5 a -AI2O3 , SiC , nuillitc . 18 9 54.6 - 36.4 1650 8 a - AI2O3 , SiC , mullite . 19 16.7 50 - 33.3 1650 8 a -AI2O3 , SiC (traces). mullite (traces). 20 90 " • 10 1775 4 SiC, AI2OC .ctAhOT (traces). 100 >• u> 90 s* •a 80 2 70 o > o 60 a, rr 50 J I 60 UI £ 50 •a £ to 4) 30 • 23pm SiC-15pm Al203wet mixing > o 5.5pmSiC-\5pm AI2O3 .. 20 0 23 pmSiC-AI2O3precipitated S? 10 ' ' ! I I I I L _L 0 10 20 30 40 50 60 70 80 90 100 mole V. AI2O3 Fig.11 hEffect of AI2O3 content on densities of green and pressurelcss sintered at 1825°C th IAl203preci- piloted }. 100 10 20 30 i.0 50 50 70 80 90 100 mole V. AI2O3 Fig.l 2):Weight loss from pressvreltss sintered AI2O3- SiC compocts al 1825 °C . th IAI2O3precipitated 1 -522- 16O0 1200 h •1200h <1 1600 h - 2000h - 2t00(- -2800 - 3200 300 700 1100 1500 1900 1300 1700 Temperature, °K Fig.( 3) Free energy-temperattjre diagram for some selected reac tions . I7 N This diagram is based on the idea thot SiC is solid. -523- 600 300 700 . 1100 1500 1900 2300 2700 Temp,°K Fig U) Free energy-temperature diagramtor som e selected reactions. F. N.This diogram is based on the ideo that SiC is vopor. -524- SiOi + Al202(g) = SiO(g) f AI2O3 <3) were negative to lenipcratures higlici than 2700 °K . This implies . (a) the activity of AI2O3 and SiO^j,) is much higher llian the other sides . (b) AI1O3 is more slable than SiCh and (c) the tie line AI2O3 - SiO exists in the system Si - Al - O Obtaining a weight loss higher than that observed for pure SiC and pure AhOj indicates that chemical reactions occur between them causing the formation of volatile compounds , Fig.(3) Careful examination of (he powder surrounding the specimen indicates no change in chemical constituents which indicates that a gaseous product escaped . Some shiny phases and needle structures were developed on the carbon plugs and on the cold pans of the heating element. X-ray diffraction showed that these phases consist of noncrystalline SiOi . 8 and a - AHO3 , sillimanite . tridymite. AlgSiCy , AJ4S12C5 , AHOC . and A^SiC^j. These are the products of reactions of the vapor phase with the furnace constituents . The effect of time.( for isothermal treatments ), and the effect of temperature, ( for constant heating periods ) , on weight loss were studied for selected mixtures and are shown in Figures (5) and (6). The following reactions were studied to assess the possible gases which can evolve when SiC and AJ2O3 are heated together: 2AJ + 2°2(g) SiO(g)' + CO(g) = SiC + 2Ah03 (4) 4AIO(g) + SiO(g) + CO(s) = SiC + 2AI203 (5) A1 + 2°(g) SiO(g) + CO(g) = SiC + A1203 (6) 4A1 SiC + °2(g) = 2A1;03 + CO(g) + SiO(g) (7) From Fig. (3) , which shows the variation of standard free energy change against temperature it can be concluded that above 1250 °K the gas phase consists mainly of SiO. CO. and AIiO: this is supported by Jackson et. al. and Misra results (8.9). To determine the type of reactions occurring, mixtures of SiC-AI2O3 and SiC - mulliie were prepared and heated at the temperatures and the periods indicated in Table (I) . Residual phases existing in the mixtures are also shown in the table . The presence of the starting phases show that equilibrium was not attained in all the mixtures but indicates the phases compatible with those lost as gas . The existence of AI4O4C, AI2OC, and AI4C3 indicates that the residual phases are getting poorer in oxygen . Silicon carbide was found to be depleted. (in some specimens disappears or residual traces are found). This means that the main constituents of the gas are SiO and CO . The compatible phases were used to construct the tie tetrahedra in the quaternary system Si - C - Al -Oat high temperatures. Fig. (7) .The compatible tetrahedra in the system are : 1 - SiO - SiO: - CO - Mullite . - - AI2O3 - SiO - CO - AI4O4C 3 - SiO - AI4O4C - CO-AhO. 4 - SiO - CO - AbO - Al2OC . 5 - SiO - CO - AhOC - SiC . 6 - AI1O3 - AI4O4C - SiO - AhO. -525- :ig. 1 5): Effect of time on final relative density end weight loss. Pressureless sintering. -526- 7-SiO- AI-.O - AhOC - SiC . The last two tetrahedra do not intersect the section SiC - Si02 - AhOj - AI4O3 . The intersection of the five tetrahedra with the section gives the triangles shown in Fig. (8). This diagram is not planar and consists of a projection of the triangles produced by intersecting the compatible tetrahedra . These tetrahedra correspond to the following chemical reactions: 4Al203 - 3SiC = 2AI4O4C + 3SiO(g) + CO(g) (8) 3Aip3 * 3SiC = AI4O4C + AJ2OC + 3SiO(g) + CO(g) (9) 2Aip3 - 3SiC = 2AI2OC + SSiO^) - CO(g) (10) 3AI203 * 6SiC = AI2OC + 6SiO(g) + AI4C3 + 2CO(g) (l«) 2AJ1O3 - 4.5SiC = AI4C3 + 4.35X3^) * I -5CO(g) (12) Ignoring the volatilization losses due to formation of Ai2C>3,the above reactions correspond toweight losses of 30.3 . 37.6 .49.4. 58.7 and 62.5 % respectively/These losses were plotted in Fig (10) against the initial molar compositions. The curve obtained represents equilibrium weight loss and is evident that it is similar to that obtained experimentally at 1825 °C for mixtures firedfo r 4 h . Fig (2). The changes in density occurring during firing can be explained in view of the above results and conclusions . For higher firing temperature, increasing the temperature increases the weight loss and the gases created can escape through the open pores . The high temperatures enhance diffusion and sintering. The alumina content increases due to loss of SiO and CO as gases and increases the sinterability. This behaviour is shown for the mixture containing initially 50 % AbOj when fired up to 1775 °C. Fig. (9). Weight losses increase exponentially with temperature, Fig. (6), causing rapid formation of gases . These gases cause drop in the density due to bloating and cracking of the specimens during the isothermal treatment . However . the initial loss of gases and their escape through the open pores causes an increase in the alumina content which facilitates sintering allowing densities above 93 % theoritical density to be achieved in mixtures initially containing 70 - 85 % AI2O3 . Further heating and production of gases difficult to diffuse caused a general decrease in density due to cracking which was observed in microstructure . Mixtures containing initially SiC and mullite showed loss in weight higher than the corresponding mixtures containing SiC and AI0O3 . Firing and equilibrium composition of each for Ih at 1775 °C . caused a weight loss of 45.5 % in case of mullite and 11.4% in case of AI2O3 . Using mixtures having 14 % mole mullite and 14 % mole Al203 and firing for 0.5h at 1825 °C caused 22.4 % weight loss in case of ntullite and only 5.4 % weight loss in case of AI2O3 . Mixtures rich in mullite. ( 14 and 15. Table I ) , showed that after firing , alumina was developed at temperatures as low as 1675 °C after lh. This indicates that the following reaction occurs: 3AhO3.2Si02 + SiC = 3AI203 + 3SiO(g) + CO(g) (13) -527- /u o 10 •/• Al203. 0.5h 60 - a 25V. Al2O3.0.Sh f • 50 V, AI2O3. I.Oh 50 - f / 2 ^0 / w p o 30 / "S. 20 { I * 10 l , J 0 1 ' 0 1£00 r 150-r0 , 18O, 0 2000 2200 Temp- . I°C1 Fig. I 6 ): Effect of tpmperoture on weignt loss by pressureetess sintering. Fig I *1« s*_ C - At* 0 quoternory system. -528- M tMollit* S *S.02 S' rSiO A sAl203 Si content etoms 3S;02 2Al203 ^•r-r .cot 3SiO Alt04C • CO -< 2AIJ0 c 1.33 CO c o u c 2A120C 9 >» X o AI4C3 3 SiC 0 0 0 01 ij r» 0 <*» 10 ••• *•* ••« > > > «» *- CD in V> M n •r0 > O *» in •" *••" Fig.(8) SiC- SiC>2 - AI2O3 - AlAC3 quaternary section F N The numbers shown here in Fig. 8. correspond to the compositions given in Table ID. -529- 100 o 10% AI O3.0.5h 90 2 ° 25 V. At2O3.0.5h • 50 % Alj0 ,1.0h c 80 h 3 •a Of 70 > 50 az 50 40 _L _L JL L_ X 0 1400 1600 1800 2000 2200 7emp.|eC J Fig.19 ): Effect of temperature on the final relative density by pressureless sintering. VI o '5 s 40 50 60 70 80 mole % Al203 Fig. (10): Calculated weight loss -530- 10 15 20 25 Pressure IKpsii 10 1S 20 25 Pressure 1 K psi 1 Fig.Hl) Effect of Hipping on the relative density and weight loss of 5SV. Al203- 45% SIC 182S"C and O.Sh. -531- The reaction indicates the existence of an extra conjugation tetrahedron , (Mullite - Alumina - SiO - CO). which is added to Figures (7) and (8). Mixtures richer in SiC than that shown by this reaction will lie in one of the other previously established tetrahedra . Accordingly. the following series of reactions occur: 3Al203.2Si02 + 2.5 SiC = A1203 + AI404C + 4.5 SiO(g) + l.5CO(g) (14) 2( 3AI203.2Si02 ) + 6.5SiC= 3AI404C+ 10.5SiO(g) + 3.5CO(g) (15) 3AI203 2Si02 + 4SiC = A1404C + AJ2OC + 6SiO(g) + 2CO(g) (16) 3Al203.2Si02 + 5.5SiC = 3Al2OC + 7.5Si0(g) + 2.5CO(g) (17) 3Al203.2SiO: + 7SiC = Al2OC + A14C3 f 9SiO(g) + 3CO(g) (18) 2(3AJ203.2Si02)+15.5SiC= 3AI4C3 + 19.5SiO(g) + 6.5CO(g) (19) Mixtures shown in Table (1), are in harmony with these reactions. It should be noted that trials to prepare mixtures of SiC and mullite by firing SiC, Si02 and A1203 caused the rapid disappearance of Si02 and SiC , [ mixture 9 . Table (1) ], due to the formation of SiO and CO gases according to the reactions : 3Si02+2SiC = Si(>) + 4SiO(g) + 2CO(g) (20) 2Si02 + SiC = 3SiO(g) + CO(g) (21) Calculating the free energy change for these reactions showed that reaction (21) is more favourable thermodynamically. Since all the above reactions produce lot of gaseous products , they can be retarded by using high pressure during firing . Accordingly, mixtures of SiC - A1203 and SiC - mullite were fired in the hot isostatic press (HIP), and the reactions were effectively retarded. Using different pressures in the range from 5K PSI up to 28K PSI , Fig. (11), for half an hour at I825°C lowered the weight losses for the mixtures containing initially 55 % mole AI203 , from about 20 % weight at about 2 atmosphere to about 3.5 % weight at 28K PSI. The high density observed at pressures lower than 7K PSI can be attributed to losses due to gas formation and encircling the mixtures with Al203 , thus facilitating densification . Since pressures above I OK PSI were used by gradual applying pressure during heating, slight densification was observed at higher pressure. For SiC - mullite system the pressure was more effective. At 1825°C a mixture containing 78 26 % mole mullite lost 22.4 % weight after 0.5h and on increasing the pressure to 28K PSI only 0.9 % weight loss was observed . Another mixture of SiC and mullite heated for lh at 1775 °C caused a loss in weight of 45 % which was cut to only 0.6 % at a pressure of 28K PSI . Accordingly , increasing the pressure prevented volatilization and in this case the relationships shown solid in Fig. (8) will prevail . It was observed experimentally that high pressure produces homogenous samples whereas normal pressure causes coring effects and rougher surfaces . Specimens containing initially 15 and 30 % mole Ah03 fired at 1825 °C for 4h and examined by X-ray diffraction showed that segregation occurred . The surface was found to be richer in Al203 and very poor in AI404C . This may be due to the fact that the AI404C is liquid at the sintering temperature . which may have been decomposed and left the surface as -532- gaseous products . The X-ray patterns showed also that most oftheSiC reacted with AI1O3 . The summarized results on substituting alumina by mullite are given in Table (2). Table (2) Summarized results on substituting AI2O3 by mullite Composition. mole % Temp. Time Wt. loss % "C h SiC Al203 Si02 Pressureless Hipping.28KPS I 86 14 0.0 1825 0.5 5.4 1.3 7.85 54.69 36.46 IS25 0.5 5.4 0.9 SO 50 0.0 1775 1 11.4 2.1 17 50 33 1775 1 45 06 4. CONCLUSIONS: 1 - As the percent addition of fine alumina to coarse SiC increases the percent relative density of SiC - AI2O3 compacts increas. 2 - Precipitation of AI2O3 on SiC causes the green density to decrease due to coating SiC instead of filling the voids 3-Increasing the AI2O3 content leads to an increase in the fired density, this is due to the high sinterability of AI1O3 . than that for SiC . 4 - A weight loss after sinterig SiC - AI2O3 compacts at high temperature was obtained due to the formation of volatile gaseous products which are produced due to the reaction between AI2O3 and SiC . 5 - The weight loss in AI2O3 • SiC compacts increases as the alumina content increases up to 30 % mole AHO3 . or by increasing the temperature or time of sintering. 6 - Aluminum oxycarbides and carbides were detected by X-ray as reaction products 7 - Some reaction products are gaseous such as SiO, AhO. and CO, these products escape from the compacts, react with the furnace constituents forming shiny phases and needle like structures. 5 - The weight loss in mixtures containing SiC and mullite is higher than the corresponding mixtures containing SiC and AI2O3 . This is due to the high reactivity of SiCh . in the mullite, with SiC to form SiO. CO. and AI2O3 . The latter reacts again with any remained SiC to form AhO. SiO, and CO. 9- The use of hot isostatic pressure. (Hipping), during sintering reduces the weight losses especially in the case of SiC - mullite. 10 - Hipping produces homogenous compacts and pressureless sinlering causes REFERENCES : (1) Asoke C.D.Chaklader, Sankar Das Gupta, Edmond C.Y.Lin . and Boris Gutowski , "AI2O3 - SiC Composites from AluminosilicatePrecurs" . J.Am.Ceram. Soc., 75 , [812283-85(1992) (2) M.M.Dobson, "Silicon Carbide Alloys", Parthenon Press . Carnforth . Lancashire , Engl and. pp.5 .(1986). (3) W.J.Lacky , D.P.Stinton , G.A.Cerny, A.C.Schaffhauser. and L.L.Fehrenbacher , "Ceramic Coatings for Advanced Heat Engines-A Review and Projection", Advanced Ceramic Materials .2,(1] 24-30 (1987). (4) George C.Wei, and Paul F.Becher. "Development of SiC - Whisker - Reinforced Ceramics", Am.Cerant. Soc. Bull.. 64, [2] 298-304 (1985). (5)W.D. Kingery. H.K. Bowen , and D.R. Uhlmann. "Introduction to Ceramics", Second Edition . John Wiley & Sons , N.Y. . London , Sydney . Toronto. (1976) . (6) Hiroshi Kubo. Hidehiro Endo, and Kiyoshi Sugita, "Sintering Behaviour of Ultra Fine Alumina- Coated Silicon Carbide", Horizons of Powder Metallurgy Part II .Proceedings of the 1986 International Powder Metallurgy Conf. and Exhibition, Metallurgy, (1986), July 7-11, Dusseldorf. W.G. (7)S. Sundaresan . S. Kanzaki , and I. A. Aksay , "Processing of Mullite Matrix Composite Reinforced with Silicon Nitride Whiskers'', Am. Institute of Chem. Eng. (1989). Annual Meeting and Ammonia Symposium. (8) Nathan S. Jacobson , Kang N. Lee , and Dennis S. Fox , "Reactions of Silicon Carbide and Silicon (IV) Oxide at Elevated Temperatures", J. Am. Ceram. Soc., 75 ,[6] 1603- 11(1992) (9) Ajay K . Misra , "Thermochemical Analysis of the Silicon Carbide - Alumina Reaction with Reference to Liquid - Phase Sintering of Silicon Carbide", J. Am. Ceram. Soc. , 74 , [2] 345-51 (1991). (10) T.Barret Jackson , Andrew C.Hurford, Susan L.Burner, and Raymond A.Cutler , "SiC - Based Ceramics with Improved Strength" , Ceramics Transactions , Vol. 2 . Edited by J.D.Crawley, and C.E. Semler. The Am. Ceram. Soc. Westerville . OH. ,(1986). (ll)Bulent E. Yoldas , and Deborah P. Partlow. "Formation of Mullite and other Alumina Based Ceramics via Hydrolytic Polycondensation of Alkoxides and Resultant Ultra - and Microstructural Effects" . J. of Materials Science , 23 , 1895 -1900(1988). (12) K.Okada, Y.Hoslu . and N.Otsuka. "Formation Reaction of Mullite from SiO: - AUO, xerogels " J. of Materials Science Letters .5.1315-1318 (1986). (13) Kiyoslii Okada. and Nozomu Otsuka. "Characterization of the Spinel Phase from SiO: - AUOj xerogels and the Formation Process of Mullite" , J. Am. Ceram. Soc. .69.[91652-56(I986) •5*4- EG fSo/i5JS First International Spring School & Symposium on Advances in Materials Science (SAMS 94) 15-20 March 1994 Stable/Metastable Ce-TZP, Effect of Chemical Composition and Preparation Method S.El-Houte Metallurgy Dept., NRC, Atomic Energy Authority Cairo, Egypt ABSTRACT Ceria doped tetragonal zirconia polycrystals Ce-TZP could be prepared with a high percent of metastable tetragonal phase. The material prepared by the coprecipitation technique is of higher toughness and more sinterability than the materials prepared by the mechanical mixing method. A toughness up to 31 MPaVHTwas obtained for the sintered samples made from the coprecipitated powders, containing 12 moleV ceria and fired at 1500°C. The fracture surface of these samples contained about 90% monoclinic phase, while the material prepared from the coprecipitated powder of 17.4 mole% CeO, contains only a stable tetragonal phase and has a KIC of 4.6 MPavfinT. Four point bend strength from 390 up to 560 MPa have been obtained depending on the composition. Compacts made from the hydrothermally treated powders had to be calcined in order to be sintered to high densities. Differences in the properties of the materials prepared from commercial powders depend largely on the additives of the trivalent rare earth oxides as well as the ceria content itself. I- INTRODUCTION: Tetragonal zirconia (TZP) ceramics, usually prepared by doping zirconium oxide; either with yttria (to give Y-TZP) or ceria (to give Ce-TZP), are in fact very interesting ceramic materials. This is attributed to their remarkable high strength and toughness (1,2,3). Moreover, Ce-TZP may exhibit shape-memory like behavior (4,5) in addition to transformation plasticity (6,7) . Ce-TZP ceramic materials are fabricated with a high degree of metastability (8) . It was found also to be more stable than Y- TZP upon low temperature aging (9), and in moist environments (10). Only at a relatively high temperature (1200*0 in a hydrogen reducing atmosphere, Ce-TZP may be reduced and ZrCe,0, -535- phase may be separated (11). Ce-TZP exists in a single tetragonal phase over an extensive composition range 12-20 mole %Ce02, while for Y-TZP, this phase exists only over a narrow composition range, from 2-3 mole % Y203. This in fact, gives more room for further investigation and studies on such an interesting material. The physical and mechanical properties of this interesting ceramic material depends highly upon the composition, grain size, impurity content and the method of preparation. These factors have been subject to numerous investigations (1-10). More work is still going on and a lot has been published recently on the properties of Ce-TZP ceramics (13-19). The discrepancies in the published results regarding the mechanical properties of Ce-TZP might be due to the presence of some impurities in the starting materials (2,20), test methods (21), inaccuracy in specifying the chemical composition or heterogeneity. In a previous work (22) the preparation and properties of Ce-TZP ceramics containing 12 mole% ceria by the coprecipitation technique, as well as the different steps of sample preparation and testing have been given. In the present work different compositions of Ce-TZP powder have been prepared using also the coprecipitation technique, in addition to some powders prepared by mechanically mixing the starting Zr02 & Ce03 powders in a ball mill. This was to study in more detail the effect of composition and preparation method on the final mechanical properties of the sintered samples. A description of the microstructure, phases and mechanical properties is also given. 2- EXPERIMENTAL TECHNIQUE: Zirconium oxychloride from Fluka and cerous nitrate from Rare Earth Corporation, were used as starting materials, for preparing powders containing 5, 10, 12, 14, 17.4 mole* of Ce02. The coprecipitation technique was used - as described in a previous publication (22) . A 0. 2M solution of the desired composition of the mixed salts was prepared . The solution was then dripped in a vigorously stirred ammonia solution, which was kept at a constant pH of 9.0. The precipitated gel, thus obtained was washed with distilled water several times, then with ethyl alcohol. Afterwards it was dried and calcined at 600 C. The calcined powder was then ball milled and sieved. Part of the coprecipitated gel of the powder containing 12 mole% ceria has been crystallized hydrothermally in a teflon container at 200 C; then washed and divided into two parts. One of these was oven dried at 80 C and the other was spray dried. Another batch of powder containing 12 mole% CeO, has been prepared by wet mechanical mixing for 20 hours in a ball mill using zirconia balls, ZrO, from Prolabo and Ce02 99.9% purity, -536- as well as commercially available zirconia powders containing 12 mole* CeOj (TZl2Ce) from Tosoh Japan. The later powder had the chemical analysis as given by the suppliers listed in Table (l). It was also wet milled in alcohol for 20 hrs. before pressing. Rectangular bars from different powders were prepared by pressing in a rectangular steel die of dimensions (35 x S mm2) at a pressure of 100 MPa. These bars had green densities of 40% TD. The samples were then pressureless sintered in air at a temperature of 1400-1600 °C. The sintered densities were measured using the Archimedes method in water at room temperature. X-ray diffraction analysis was used for quantitative phase determination (23). A Phillips and XD-3A Shimad2U machine with Cu Kc*/ Nickel filter had been used. The samples were ground and diamond polished using Abramine grinding polishing machine from Struers. They were thermally etched at 1350 * C for 20 minutes, and the microstructure examination was done using a scanning electron microscope J5M-T20 Jeol. The hardness and fracture toughness -using the indentation fracture technique (IP)- were measured on the polished sample surfaces, using an Otto Wo 1 pert Werke Hardness Testor machine with loads ranging from 200-1250 N. The measurements were taken at each load, using an optical microscope BH2-Olympus-Japan- Niihara's equations (24) were used for the fracture toughness calculations. A four point bending with 10/20 mm loading/supporting spans and 0.22 mm/min cross head speed was used for fracture strength determination. Pour samples were tested for each measurement. Table I Chemical Analysis of TZl2Ce in wt'.fc "from Tosoh Corp. Tokyo, Japan" Ce03 Si02 Fe20, Na20 16.55 0.002 0.005 0.0119 RESULTS AND DISCUSSIONS: a-Effect of Compositions: The compacts made from powders prepared by the chemical coprecipitation technique showed excellent sintering behavior (22) . Densities over 98% of the theoretical densities wer« obtained for all the samples of different ceria compositions, except those of ceria content less than 12 mole%. The XRD patterns -537- obtained for the as-sintered sample surfaces showed a single tetragonal phase for the samples containing 12,14,17.4 mole* ceria, as shown in Pig. 1. Samples containing S mole% ceria showed only a tnonoclinic phase, while samples containing 10 mole* ceria showed monoclinic and tetragonal phases. The 5 moleV ceria samples contained cracks which were produced during cooling as a result of tetragonal to monoclinic transformation which renders their mechanical strength very weak. The samples containing 10 mole* ceria were apparently sound with no visible cracks, but with relatively low densities of about 87.8*TD of the tetragonal phase. Consequently, they had very low strength. So, the 5,10 mo lev ceria materials were excluded from further investigation. The micrestructures of the samples containing 12,14 and 17.4 mole* ceria are shown on the SEM photographs, Fig.2. It could be seen that the material containing 17.4 mole* ceria has relatively smaller grain size than those of the other two compositions. Fig.3. shows the variation of the mechanical properties: Vickers hardness, fracture toughness KIC and bend strength, as functions of composition for the samples fired at 1500 C in air for 1 hour. It can be seen from the figure that a maximum in the bend strength; (560 MPa) corresponds to the 1-3 mole* ceria material sintered at 1500*C. The fracture toughness KIC was found to decrease with the increase in ceria content, while the Vickers hardness increases. This is in agreement with published data on Ce-TZP (1). On the other-hand the fracture toughness KIC for the composition 12 mole* ceria has been reported to increase with sintering temperatures, while the Vickers hardness decreased (22) . Taking in consideration that fracture toughness values generally depend on the test method (1,2.21,22) and on the equation used for calculation (22), a KIC of 41.1 MPa Fig.5 shows optical micrographs of Vickers indentor impressions on the polished surfaces of the sintered materials, containing 12 and 17.4 mole* ceria. A rosette shear-like shape could be seen in the first as a result of the stress induced transformation of the -538- c- 17.4 sola % evrli m m J { JUL_ TlgA XID H*t«r* »f th* ct alater*4 murtmct ••«• froa p, eopnclflut«4 po«d«r>. I Fig.2 SEM micrograph of the thermally etched samples containing: a) 12 mole* ceria, b)14 mole*, c) 17.-5 mV. -539- ric.J VlcXr'a kartfmaa. trmctur* (H|UtH aa« atM •trraeth vrraaa caasaaltlaa ror aaKrlaU alnUrH al 1500«C far Ik. V»W*w**fc^ii T—i—r rit.4 MB aattara aa«a aa taa fractvrt aarfact at tka •atarlalat «,V.«. (tatatnfsaal. auaonMllalc). Ftf.S SaaaraKl cantraat aptlcal «lcr«sr»»h* «r tht ln- tftata aaa> aft tka aollakaa aurfacaa aft a) Miarlal •/ Maala 'A ccrla •lntcrrd at 1(00*C/lb lB«nlat|aa laa« Uitm. *0«. ») aalarlal at IT.4 moUT- «»rla .tntrrri al 16Cf*r/)a. Itrfaatatlaa laa« 2CC5. Mi. -540- metastable tetragonal phase, while the latter shows spalling and long cracks due to the complete absence of transformation toughening in the 17.4 mole% ceria samples- as explained before. The indentation loads are given in the figure captions for comparison. The XRD pattern Fig.6 gives more evidence about the transformability of the 12 moleV CeO; material and the stability of the 17.4 mole% CeO,. They are measured on the cut surfaces of these two samples after diamond cutting. In the pattern corresponding to the 12 mole% CeO,, the monoclinic phase which results from the transformation of the metastable tetragonal phase could be seen, while for the 17.4 mole* Ce02 an increase in the intensity of the tetragonal (002) peak at the expense of the (200) one was observed. A similar XRD pattern obtained for the ground sample surface of this material is given in Fig.6-C. This is referred to as a domain switching or reorientation, which indicates that the stable, tetragonal zirconia is a ferroelastic material and a new toughening mechanism could be added (26-28) to the transformation toughening i.e the ferroelastic domain switching mechanism. Microstructure and X-ray diffraction evidences are in support of the occurrence of domain switching (29-31). In the present work, the XRD patterns obtained for the cut and ground sample surfaces of the 17.4 mole% ceria gave additional evidence, where the exchange of the (002) and (200) peak intensities occurs, without any trace of the monoclinic phase. b- Effect of the method of preparation 1- Mechanically mixed and coprecipitated powder A comparison between the properties of the samples prepared in the conventional way, by mechanical wet mixing the powders, and those prepared from the coprecipitated powders, will be helpful in this respect. The powder of composition 12 mole* ceria will be taken as an example. The densification behavior of the mechanically mixed compacts (MM), and those from the coprecipitated powder (CP) are shown in Fig.7. It is clear from the figure that the (CP) samples are more sinterable than the (MM) samples, specially at low firing temperatures. Fig.8 shows the XRD pattern of the material fired at 1650 C. It contains about 7% monoclinic phase. Fig.9 shows the SEM micrograph of the microstructure of an (MM) sample sintered at 1500°C for 1 h. It has relatively small grains compared to the coprecipitated material; Fig.2. The mechanical properties of this sample were; bend strength of 4 04 MPa, Vickers hardness of 7.8 GPa and a toughness of l2MPa\Tm". These values are superior to those published for the materials prepared by the same method (32,33). As for the fracture toughness, it was lower than the value obtained for the (CP) material sintered at the same temperature. This may be explained by the relative increase in the concentration of the ceria in the -541- near surface grains than that in the bulk. This could be considered as microinhomogeneity compared to the macroheterogeneity expected to exist in that material of 18 mole* ceria which contained a monoclinic phase after firing (25) . 2- Hydrothermally treated powder Compacts made from the hydrothermally crystallized powder either through freeze or spray drying, had low sintered densities of 784 and 80V TD of the tetragonal phase respectively. 70* monoclinic phase was found on the XRD pattern of the sintered sample surface as shown in Fig.10. Some differential sintering might have taken place due to the presence of agglomerates (34). resulting in cracks upon firing (35) and consequently a weak material. It was found that calcination of this powder at 600 C for one hour followed by crushing and sieving might produce a material of 99VTD upon firing at 1500 C for one hour. This, in fact, put some doubt about the validity of the hydrothermal treatment of powders for producing strong and dense ceramics. 3- Commercial powders Compacts of TZl2Ce made by directly pressing commercial powder from Tosoh-Japan, taken directly from the bottle gave low sintered densities 76VTD after firing for l h at 1500 C. This was observed before for similar powders (3,36). However, when this powder was wet milled in ethyl alcohol before pressing, a sintered density of 99VTD could be obtained. The sintered material was of 100V tetragonal phase. The microstructure is shown in Fig.11, which shows a fine microstructure with smaller grain size than that of the coprecipitated material (12 mole* ceria-Fig.2). The fracture strength and fracture toughness were 630MPa and 13.2 MPa -542- . CP 9: f MM" 9o / 9^ / J? [ 90 1400 1)00 1600 b-cut surface of f a mt • lntann/: ttaa.'c ceria. Mc.7 Dmai Heat tar bafcatriar al caapwla •*>» fraa saehanlaally ataat ra*a>r. Caspar** la tbaaa fraa caareeipiiataa' aaatfara. w".8 _ rif.i 1»D aaltanu af aattrlala alataraa at 1*00*C. flt.8 ZXO at tba aa alatara* aaapla aurfacaa aalt fraa aachulcally atxta s*a*ara. Fig.9 SOI alerograph af th« atchanlcall? aliad paadar • Uttria at 1}00*C. -543- —,—. —r 61] ** i a o o £* « J UUl AM:1 > » » n >i jj^ ii * » a >r^ Fig.10 MB pattern of a sintered canple surface at 1500 C. made free) hydrothermally treated powder. Pig.11 SEM micrograph of TZ12Ce sintered saayle at ISOO C. -544- CONCLUSIONS: 1- Using the coprecipitation technique, homogeneous materials with good mechanical properties i.e toughness and 3trength could be prepared. 2- The sinterability of the powder prepared by the coprecipitation technique is higher than that prepared by mechanical mixing. 3- The metastability of the tetragonal phase as measured by the amount of the monoclinic phase produced on the fracture surface, is a decreasing function of ceria content in the range 12-17.4 male* caria. The material of 12 moleV ceria contains mainly a metastable transformable tetragonal phase and a KIC of 41.1MPa m could be obtained, while the 17.4 mole* ceria material had only a stable tetragonal phase and a low fracture toughness of 4.6MPa m. 4-Compacts from hydrothermally treated powders showed low densification and a monoclinic phase on the sintered surface. 5- The microstructure examination of the sintered materials containing 12 mole* ceria showed smaller grain size for materials made from the commercial powder (TOSOH) and from the mechanically mixed powder relative to those made from the coprecipitated powder. 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Editors, J.J.Bentzen, J.B.Bilde Sorensen, N.Christiansen, A.Horsewe11 and B.Ralph, Riso, Denmark.1990. 23- . H.Toraya, M.Yoshimura and S.Somiya, "Calibration for Quantitative Determination of the Monoclinic-Tetragonal ZrO, System by X-Ray Diffraction", J.Am.Ceram.Soc., 67,cll9-cl21. 1984. 24- K.Niihara, R.Morena and D.Hasselman, "Evaluation of KIC of Brittle Solids by the Indentation Method with Low Crack-to-Indent Ratios", J.Mater.Sci.Lett.. 1,13-16, 1982. 25- T.W.Coyle. W.S.Coblenz and B.A.Bender, "Transformation Toughening in Large-Grain-Size CeO, -Doped ZrO, Polycrystals", J.Am.Ceram.Soc., 71{2)C-88-C-92. 1988. -546- 26- A.N.Virkar,"Ferroelastic Domain Switching as a Toughening Mechanism in Tetragonal Zirconia", J . Am.Ceram.Soc . ,69(10) , C-27.4-C- 226, 1986. 27- A.N.Virkar,"Toughening Mechanism in Tetragonal Zirconia Polycrystalline (TZP) Ceramics", Advances in Ceramics, vol.24: Science and Technology of Zirconia III. 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Symposium on Advances in Materials Science (SAMS 9-1) 15-20 March 1994 RADIATION DAMAGE OP METALS BY HIGH-ENERGY CHARGED PARTICLES V.F.Reulov Joint Institute for Nuclear Research, Flerov Laboratory of Nuclear Reactions, 1419S0, Dublin, Moscow region, Russia The methodological aspects of irradiation of IIKLCP materials arc developed of alpha* particle and proton irradiation; calculations of PKA spectra and experimental data on accu mulation of transmutation gaseous impurities (helium and hydrogen) and methods of inform Itclium (hydrogen) doping over the volume of samples and also formation in the samples of doping profiles with various space orientation. The experimental data on defect structure formation in molybdenum depending on energy and dose of proton and alpha-particle irra diation are presented. The problems of influence of PKA spectrum and impurity helium on molybdenum radiation hardening arc studies. INTRODUCTION High-energy light charged particles (HELCP), alpha-particles ard protons with 10-50 McV energy, for example, arc known to be considered in the capacity of bombarding particles sim ulating the effect of neutrons of fusion reactor in the conditions of the absence of operating fusion reactors and high-flux sources of 14 McV neutrons [1,2]. However, practical use of IIELCP is an extremely complex problem since at present there arc some difficulties in the experiment performance on high-energy accelerators, cyclotrons in particular. The methodological aspects of irradiation by HELCP including new developments of irradi ation dosimetry and calculations of total energetic PKA spectra, the methods of uniform ion implantation of bulk specimens, ways of ion implantation profdc formation in the samples with various spatial orientation arc treated in the paper. The experimental data on molybde num defect structure and radiation strengthening depending on bombarding particle species and energy and helium concentration arc also presented. CALCULATION OF THE' TOTAL PKA SPECTRA flic energetic PKA spectrum plays a determining role at the stage of generation and devel opment of displacement "cascades in the irradiated material so further fate of the radiation defects, their temperature stability, tendency for clustering, etc. depend on it. To calculate a total number of displacements it is necessary to obtain first of all the energetic I'KA spectra from elastic and inelastic interactions of IIELCP with the target atoms, then to determine damage energy and choose model of forma lion of secondary displacements on the basis of any model presentation. Energetic PKA spectrum from clastic .scattering is connected with dilfm-nlial miss*stvLion of elastic scattering. To obtain inelastic I'KA spectrum one should know integral iToss-secliim of all inelastic processes. -549- Calculations of inelastic and clastic cross-sections of I1ELCP interaction with the target atoms have been performed using the methods of parameterized phase analysis (PPA). Dy means of the code based on PPA method the differential cross-sections of elastic scattering and integral cross-sections of inelastic scattering for a scries of elements in a broad range of IIKLCP energies have been obtained. On the basis of the mentioned approach the code "DAMAGE" to calculate total FKA spectra for a wide class of materials under proton and alpha-particle irradiation in 1-5-50 MeV energy range has been developed |3). Figure 1 presents total energetic PKA spectra in molybdenum under alpha-particle and proton irradiation. PKA Energy, keV Figure 1. PKA energetic spectra in molybdenum irradiated by alpha-particles (a) and protons (b). Alongside the radiation defect formation in materials irradiated by HELCP the accu mulation of nuclear reaction products (NRP) lakes place. Plotting of the spectra of NRP masses docs not present any difficulty if there arc relative data on reaction cross-sections. Upto-now, the greatest practical interest of all the spectra of NRP masses is devoted to the region containing hydrogen and helium. In this connection the experimental data on helium accumulation in a scries of metals under proton irradiation and hydrogen accumulation un der alpha-particle irradiation have been obtained (Figure 2) [l,5|. On the basis of radiation damage correlation made on energetic PKA spectrum and NRP mass spectrum, helium and hydrogen in particular, it has been shown that alpha-particles with 304-50 McV energy and protons with 174-30 MeV energy can serve for simulation of fusion neutron effect on materials. -550- Ic E I 2 Energy, MeV Figure 2a. Helium formation in natural Cr and Fc(I), Nl(2), Nb(3), »\fo(«l), Zr(5). 2b. Hydrogen formation in natural molybdenum (I), zirconium (2) and niobium (3). METHODS OF IRRADIATION AT CYCLOTRON With the aim of developing a method of dosimetry of JIELCP irradiation combining high- sensitive and a broad energy range a now material has been evaluated, nnm.ily natural multi-isotope molybdenum in which the total output of several channels of nuclear reactions of (or.xn) type, for example, of five molybdenum isotopes (*'.9*Wr'98Mo) is used. As a result of such nuclear reactions one and the same 97llu isotope is formed, whose output docs not depend on IIELCP energy, for example, alpha-part'ules in 20-rflO MeV range (Figure 3). The problems of providing the conditions of uniform ion implantation over the volume of bulk specimens by different high-energy ions have been overcome by two irradiation tech niques. The first uses atmospheric air as a filter with constantly changing energy degradation in conditions of reciprocating movement of the irradiated specimen in it. In the second, the specimen is rotated constantly in particle fluence. Calculations by the "DAMAGE" code averaged over the total thickness of the specimen of the ratio of helium concentration, for example, to the number of displacements for many materials is about 0.3-T-0.5 at. %Ile/dpa. The possibility of direct study of the defect structure forming immediately in the particle straggling 2<>ne would allow understanding of the processes of ion implantation of semicon ductors and metals and also peculiarities of heavy ion application for simulating high-dose effects after neutron irradiation of materials. However, microscopic scale of the partirle straggling zone in materials after their irradiation by light and heavy ions does not allow many traditional experimental methods of radiation physics to be used to study the radia tion effects in this '/one. 551 •.-•*%. \aiOO\ »* oc 100 s '-*—.- 10 X 30 40 50 Energy. MeV Figure 3. Experimental dependence of cross-section of *7Hn formation in natural molyb denum upon alpha-patticic energy for 40 MeV (0) and 50 MeV (A) initial energy. To solve such problems, a new approach to variation of spatial location and change of geometric sizes of radiation damage energy profiles and ion implantation in the irradiated specimen has been developed. Willi the purpose of a single viewing in TF.M of the structural changes in the total energetic profile of damage of the bombarding particles, an irradiation technique has been developed consisting of ion bombardment of the specimen through a degradating filter of cylindrical shape. The structure of molybdenum in the straggling zone of nitrogen ions with SO MeV energy after annealing at 1100°C is presented in Figure 4. With the purposes of forming in subsurface vohimc of the specimen the damage profile with the size many times (hundreds of limes) that for traditional irradiation techniques, the spec imen is irradiated by mo9no-cncrgetic charged particles through the degradating filter in the shape of a foil bent according to the prc-sct law. Tims, for the given initial energy of bombarding ions.it is possibile to generate on the bombarded surface on energy profile with a multiply-enlarged scale of losses of bombarding particle energy. It provides possibility of more detailed study of changes of those or other material proper- tics depending on bombarding ion energy, in their straggling zone particularly, by means of nuclear-physical methods not acceptable for these purposes earlier. Illustration of the use of this method to study changes of lattice parameters and microhard- ncss of molybdenum in the alpha-particle straggling zone with 29 MeV energy is shown in Figure 5. 552 I-.-. Figure A. Zone of straggling in molybdenum irradiate.! bv lle+ ions after annealine at HOCC. ' U g Figure 5. a) Calculation profiles of U 'o concentration 0,01 * a • t « Depth, mm RESULTS AND DISCUSSION Alpha-particle and proton irradiation of molvUlemmi in 10-4-50 and IO-fM McV energy range, respectively, with damage levels from 5-l0"« to -M0-" dpa at temperature not more than (iO°C causes defect formation i„ the fort,, of clusters and dnlocatton loops in a narrow range of sizes <|-K'» ,„.,). Size .h-fa-i «Iisl.il,„ti«.u and their imiiilicr density 0) .£" M « JO » to Jr e a «0 « 4 •a i • ' ~ f U » 19 IS 10 Particle energy, MeV Analysis of the type of the observed defects by means of 2-1/2-D method has shown that more than half of the observed defects arc of vacancy type, the other interstitial. In the whole range of the investigated doses, saturation of clcfcct number density is not observed and their mean size is not significantly changed. It should be noted that for the same dam age level, alpha-particles create approximately twice as many defects visible in TEM than protons. Using phosphorus ions in energy range from 5 to «10 kcV it is shown that the observed va cancy clusters arc formed athcrmally, i.e. by cascade displacements. It has been determined that threshold of PKA energy causing the formation of visible in TEM vacancy cluster of 0.7 nm size, is about 10 kcV for molybdenum [G]. In the experiments on molybdenum irradiated by fission fragments of uranium, neon ions with 137 MeV energy and neutrons causing formation of PKA spectrum with a wide energy range, it is shown that the absi-noc of vacancy clusters of more than 5 inn size is condi tioned by subcascadc processes of vacancy cluster formation |7J. From the experimental data and taking cascade efficiency equal to 0.5, the threshold energy of subcascadc processes for molybdenum has been determined to be about 75 keV. The observed constancy of defect number dencity with increase of damage level due to change of particle energy in conditioned by elective cancellation of vacancy clusters formed in dis placement cascades by mobile interstitial atoms. Most effecli vein this process arc interstitial atoms formed from the low-energy part of PKA spectrum J8j. Such a feature is also con firmed by the nature of the VV value change characterizing the relation of vacancy number in vacancy clusters visible in TEM to their number obtained from PKA spectra from proton ami alpha-particle energy (Figure 7). 554 w.io-« 25 • Figure 7. Dependence of relation of the fjuantily or the observed 20 - 1 vacancy defects lo their number, IS " obtained from spectrum, upon energy of alpha-particlcs-0 and protons-Q •0 s . •-• 1 1 1 30 40 JO JO 10 0 Ion Energy, MeV It is seen that the greatest change of W value takes place in the case of alpha-particle irradiation since the portion of point defects in total PK A spectrum increases more strongly with decrease compared with protons. The results of measuring of microhardncss along the path of alpha-particles and proton mo tion arc given in Figures S and 9. Figure 8. Dependence of •(•particle energy. MeV molybdenum microliarducss u «s«ou»ua*M0 upon energy of bombarding WO- alpha-particles for •4 1 flucnces: 210" m"a (1); E 71 3 E M 4.210 m" (2); cr» 7-10-' in"2 (3). • \ M • ^-~T1Z"--?J^ 0 M 0,1 M o> a> Nplh, mm It is seen that growth of molybdenum strengthening docs not change in alpha-particles energy interval from 10 lo 50 MeV despite of the fact that damage level during this process changes almost by the factor of 5. In contrast to alpha-particles, molybdenum irradiation with protons causes iiomnonotonoiis change of hardening growth in 10-2-30 mcV energy in terval, and in 10...15 MeV energy region the effect of radiation loss of strength is observed. As mentioned above, number density and mean size of defects visible in TEM in the same energy range is not changed. The experiments on strength growth with level of molybdenum damage irradiated by high-energy protons and alpha-particles and also fission neutrons (Figure 10) reveal an im portant peculiarity: the value of strengthening growth for the same damage level increases with decreases of stiffness of PK A energetic spectrum. 555 proton energy. MeV Figure 9. Change of molybdenum microhardncss depending on proton energy for (lucncos: 5-10" (i); 110?o in"2 (2); 510- (3); MO" m-5 (-1); 1.210" m-J (5). o o.t a.\ o.t OLS 1,0 i.i i,< i,» U 1.0 Depth, mm Figure 10. Change of molybdenum microhardncss depending on damage level for irradiation by:A-neutrons, +-protons, 0"a'P^a-Par''c'es- Damage, dpa In the experiments on molybdenum implanted with helium and irradiated by different en ergy alpha-partich.'s it has been found ['J] that the value mid nature of the change of radiation strengthening due to helium concentration also depend significantly on energy conditions of formation of defects structure, responsible for strengthening in the presence of helium (Figure 11). It is seen that the biggest growth of niicroharducss value lakes place under irradiation by low-energy alpha-particles. CONCLUSION -Code "DAMAGE" has been created to calculate total energy I'KA spectra in a wide class of materials under alpha-particle and .proton irradiation in 1...50 mcV energy range. -Helium and hydrogen accumulation has been calculated by experimentally measured exci tation functions in a wide class of materials under proton and alpha-particle irradiation in 20...50 McV range. -A new methods of specimen irradiation by high-energy charged particles has been providing: a) uniform ion implantation over the bulk specimen volume; b) change of sizes and spatial location of energy profiles of radiation damage and ion im plantation in the irradiated specimen. -556- 100 .^ S--.10 TCrTOTPTTQO » i-iu I * •" • • i • JO* I0"4 JO'3 \0'1 10*1 C„e. at.% Figure II. Dependence of degree of radiation hardening upon concentration of helium implanted for various PKA energetic spectra (©-/?„ 3 McV; (&-/?„ 15 McV; 0-B„ 29 McV). -It is experimentally proved that vacancy clusters observed in molybdenum irradiated by HELCP have been formed in cascades of atom-atomic collisions. -!t has IMJCII stated that subc;iscade pror<:ss of drfectfomialion is the reason for the observed independence of maximum size of vacancy <|tisl<-rs (.1 urn) njxjti I'KA spectrum M illness. l'KA threshold energy for formation of cascade (10 keV) and snlicascadc (T.'i keV) regions visible in TEM in molybdcnutn has been determined. •The determining role of energetic l'KA sj>ectrum in the formation of defect structure re sponsible for radiation strengthening has been found. Il has been shown that for the same damage level the greatest contribution is made by defects formed from the low-energy part of PKA spectrum. REFERENCES: 1. Logan CM.: Lawrence Livermorc Lab. Report UCItL-51'22-1.1972.15|». 2. Abdrashilov I.U., Vagin S.P., Rcutov V.F. ct al.: J.Tcchn. Pliys. 19S3.53.3C9. (in Rus- sian). 3. Abdrashitov I.U., Botvin K.V., Reutov V.F.ct al.: Report 4-SO, Inst.Nucl. riiys.of Ac.of Sciences KazSSR, Alma-Ata, 19S0. (in Russian). 4. Lcvkovskii V.N., Reutov V.F., Uotvin K.V.: Ratliat.EIT. 19St.S0.223. 5. Lcvkovskii V.N., Reutov V.F., Rolvin K.V.: Atnmnaja luicrgia.I990.G9.99. (in Russian). G. llcutov V.F., Vagin S.P.: Fiz.Tverd.Tcla (Leningrad). 19S-I.-I.1040. (in Russian). 7. Ibragitnov Sli.Sh., Reutov V.F., Vagiu S.P., I3otviii K.V.: Fix. i khimija obrabotki matc- rialov.l9S7.1.3. (in Russian). S. Ibragimov Sh.Sh., Rcutov V.F., Vagin S.P.: J.Tcchn.Plms.l9S3.G.1192. (in Rassian). 9. Ibragimov Sli.Sh., Rcutov V.F., Vagin S.P.: Voprosy Aloinnoj Nauki i 'lVkhniki, Scries: Radiation Damage Physics and Radiation Technology. I9SS.1(43).1. (in Russian). -557- First International Spring School & Symposium in Advances in Material Science (SAMS 94) Cairo, IS - 20 March 1994 SWELLING OF Fe-30%Ni ALLOY DURING Ar+ IONS IMPLANTATION Soukieh, M. and AI-Mohamad, A. Atomic Energy eommiution, P.O. Box 6091, Damascus, Syria. Abstract Transmission electron microscopy and microhardness techniques have been used to study the structure and phase changes during ion beam implantation of Fe-30%Ni alloy by 40KeV AT* ions to integral doses in the range 1015 - 1017 ion? 'cm2 at room temperature. Formation of gas bubbles in the implanted films was noticed after annealing at temperature higher than 673K. The dependence of gas bubbles size concentration, and swelling rate on ions doses was also investigated. 1. Introduction Ion beam bombardment of metallic surfaces by unsoluble ions (i.e He+, Ar+ ) induces many types of structural defects in these surfaces. The implanted ions may then aggregates at the structure defects and form gas bubbles and voids*1*. The nucleation of such defects may increases the volume of the implanted material which called irradiation swelling*2*. Many experimental work has been carried out to investigate the effect of irradiation and material properties on swelling phenomenon. However, detailed and clear picture of such phenomenon, has not been fully understood. For example, the swelling of iron alloy with 30-35% Nickel is -559- very small in comparison with other concentration of Nickel'3', and there is no clear data or study dealing with the possibility of phase transition in Fe-Ni alloy during irradiation. This paper contains complementary results of our previous paper'4! which focuses on the study of phase structural change in Fe-30%Ni implanted by heavy ion accelerated at medium energy. 2. Exerimental Electron beam technique was used to deposit the thin films (65+10nm) of Fe-30%Ni alloy at 10"4 Pa vacuum pressure on fresh polycrystal NaCl slides. During the evaporation the substrates were heated up to 500K. The deposited films were then irradiated by 40KeV Ar+ ions at room temperature with doses ranging from 1015-1017 ions/cm2 with current density of 2uA/cm2. The beam energy was chosen in such a way to be equal to the mean projected range of Ar+ ions which was equal to half of the films thickness. The prepared films were investigated by a JEM-100CX electron microscope techniques. In addition, microhardness of bulky samplesQxlxO.lcm, which were polished and annealed up to 1200K before implantation was measured). The deposited films were also annealed in situ at temperatures of 673, 873, and 1073K for one hour in 10'4 Pa vacuum pressure. The swelling ratio was calculated using the following relation'5^ : Av _ £ y i3 n r a v ~ 6 ^ \) \) where n^ is the mean bubbles density, d., is the mean bubble size. -560- 3. Results and Discussion Diffraction patterns and structure of the implanted films are shown in fig. (1). Analysis of these patterns shows that films structure consist of three polycrystalline phases : a-Fe, y-Fe, and ordered phase Fe3Ni (see Table (1)). Formation of this ordered phase (Fe3Ni) can be explained by diffusion stimulated by irradiation of complexes such as Fe-V which have high mobility. This process may cause splitting, in the structure of the irradiated material, into two ordered phases Fe3Ni and FeNi^J. Formation of gas bubbles in the implanted films also was noticed in the films structure after annealing at 673K for 60 minutes, as shown in fig. (2). After annealing the implanted films with dose (1016 ions/cm2) at 673K, Ar+ gas bubbles size was in the range of 5-40nm, and the average size was about 12nm, and their density 4.3x10 bubbles/cm3. By increasing temperature to about 873K, the average bubbles dimensions increased to about 20nm and their density increased to about 5,3xl014 bubbles/cm3. Further increasing of annealing1 temperature to about 1073K dimensions of bubbles and their concentrations started to decrease. The increases of bubbles dimensions may result from coalescence of microvoids and nucleation processes^. However, the decreases in bubbles dimensions result from dissociation and gas release of voids after reaching a critical size. Fig. (3) shows the variation of main swelling parameters; mean bubbles size d^. mean bubbles concentration n^ as a function of implanted doses. From the figure, it is clear that the bubble concentration has a pronounced peak (5.5xl014 bubbles/cm3) at an implanted dose of 5xl015 ions/cm2. The bubbles concentration and their sizes tend to saturate after implantation with doses in the range of 1016-1017 ions/cm2. -561- •Fe Fe3Nl Ffe. -562- 300 nn M.-JO,4(oi.":') M -lO^Jcm-3) 6 J||.IJ.I,._ 1(1 20 30 40 50 I.) fi Fig. (2) : Microstructre and defect size distribution of implanted Fe-30%Ni (1016 ions I cm2): a) annealed at 673 k for 60min. b) annealed at 873 k for 60min. c) annealed at 1073 k for 60min. -563- 3 -r^.lO^/cm , dv(im)t^.* 1 5 tttOO ICH5 **n» ««i/ «••# httt/tmt Fig. (3) s The dose dependence of mean bubble concentration, size, and swelling rate at 637 k. oc+oo BE*IB BCfl« tct-17 de>« Irni/oml Fig. (4) : Microstructre (AjXH I fjH) changes as a function of argon ion doses. The saturation behavior can be explained in term of sputtering effects and thermal diffusion of the unsolbule Argon atoms towards the surface and into the depth of the films. Hence, the total concentration of the implanted ions remains almost constant after implantation with dose higher than 106 ions/cm". This explanation was confirmed by the RBS data obtained in our previous work^4'. The dependence of swelling rate (AV/V) in the implanted films on bombardment doses over the range 1015-1017 ions / cm2 after annealing at 673K shows a steady increase up to the highest dose (1017 ions / cm2) as shown in Fig. (3). The dependence of microhardness of the implanted Fe-30%Ni samples on Ar+ beam doses is shown in fig. (4). A reduction of about 2% in the samples microhardness was noticed after implantation with dose of 1015 ions/cm2. This reduction may be a result of annealing stimulated by irradiation'8'. Further increasing of ion doses the microhardness was increased up to 15% at maximum dose (1017 ions/cm2). This effect is due to highly presence of irradiation defects in films structure, which may form obstacles for dislocation movement. 3. Conclusion The conclusion of this work, reported in this paper and in another publication'4', is the following : 1. The swelling parameters were found to be saturated after ions doses higher than 1016 ions/cm2. 2. The bubbles dimension and swelling ratio were increased with increasing the annealing temperature, and further increasing of the annealing temperature decreases these parameters. 3. New ordered phase Fe3Ni was formed as a result of diffusion stimulated by irradiation. 4. Microhardness of Fe-30%Ni alloy was increased about 15% at maximum dose (1017 ions/cm2). -566- ACKNOWLEDGEMENTS The authors would like to express their gratitude to Prof. I. haddad general director of Atomic Energy Commission of Syria for his continuous guidance and encouragement, and would also like to express their thanks to Prof. G. Carter (Salford University) for helpful discussion of results References 1. K.L. Merkle, Phys. Stat. Sol., 18, 173 (1966). 2. T. Muroga, H. Watanable, K. Araki and N. Yoshida, J. Nucl. Mat, 155,1290 (1983). 3. W.K. Appleby, D.W. Sandusky and U.E. Wolff, J. Nucl. Mat., 43, 213 (1972). 4. M. Soukieh and A. Al-Mohamad, (accepted and to be published in Rad. Eff and Def. in Solids). 5. T.M. Williams, J. Nucl. Mat., 59, 18 (1976). 6. P.L. Grozen and U.L. Radinov, Vaprosi Atonmi Naoky e Techniki, Seri Fiseka Radidsonikh Povrojdini, 4, 58 (1981). 7. F. Wasiliu and V. Teodoreseu, Radiation Effects, 27, 75 (1975). 8. R.A. Orlou, Physics of Radiation Effects in Crystals, Elsevier Pub. (1988). -567- First International Spring School & Symposium in Advances in Material Science (SAMS 94) Cairo, 15-20 March 1994 EVALUATION OF THE DAMAGE INTRODUCED IN MATERIALS PROPOSED FOR THE FIRST WALL OF FUSION REACTORS USING TRIM CODE PART I: NEUTRONS EFFECTS LA El-Shanshoury, G.I. Orabi, AX El-Shanshoury Department of Materials Engineering, NCNSRC, ATOMIC ENERGY AUTHORITY, Cairo, Egypt. Abstract The damage introduced in three different steels by fusion neutrons (14 MeV) was studied using the computer program for the transport of Ions in Matter (TRIM). The materials studied are type 316SS (17% Cr, 12% Ni), AMCR (10% Cr, 17% Mn), and MANET steel (10% Cr, 0.6% Ni). Neutrons effects were stimulated by those of hydrogen ions (protons) accelerated to energies of 17 MeV which produce equivalent value of displacements per atom (dpa). The paths of ions and tracks of recoils, the ion range distributions, the recoil distributions, as well as vacancy distributions were recorded as a function of target depth. The results showed that the average number of recoils and vacancies/ion (interstitials/ion) for the three steels ranged from 180 to 243 recoils and 121-137 vacancies/ion, respectively. The corresponding replacement collisions were evaluated to range from 33 to 43% of the total number of recoils. Total energy loss to recoils and ion energy to recoil were estimated to be 16.7-29.6 KeV and 8.28-10.6 KeV, respectively. The corresponding values of energies per recoil are 93-140 eV/recoil and 42-46 eV/recoil. The total energy to vacancies ranged from 3.0 to 3.43 KeV and the corresponding energy to create a vacancy is about 25 eV. The energy of an -569- ion to produce Frenkel pair ranges from 67 to 71 eV and the energy required to produce a recoil ranges from 118 to 165 eV. It was concluded that most of the damage introduced in the three steels was produced by the recoils before coming to stop. Although the general features of the damage are nearly the same, austenitic steels have presented less damage in the structure. 1. Introduction It is well known that at fusion reactor conditions both neutrons and a-particles are released according to the following reaction: 2 3 1 4 XH + XH -> Qii + 2He + 17.5 MeV Fast neutrons are emitted with 14 MeV. They induce severe damage in the first wall of fusion reactor^1'2! and consequently cause changes in material properties. These changes are ultimately related to the interaction of the intrinsic crystal defect, associated with the property under consideration, with the defects induced by irradiation. The radiation-induced defects are the resultant of two processes which occur simultaneously during irradiation^. The first is the collision process in which a large amount of .energy has been transfered to constituent atoms n the material, creating the so called primary damage. This damage is charaterised by Frankel Pairs (FP's) production, replacement sequences and non-displacive collision sequences^'. The general picture is dependent on the damage energy, damage rate and collision energy as well as primary knockon Atom (PKA) energy^3"5'. It is also affected by the material properties that determine how the energetic particles interact with atoms and how defects interact within the material. So the general picture of primary damage is a function of material type and irradiation environment. It -570- provides information which can be used as a base for comparing the performance of different structural materials in fusion reactor conditions. The second process is the reaction of the primary damage which develops the final defect structure. This occurs mainly under the effect of temperature where diffusion and recombination processes operate effectively. Another importance of studing the naturae of primary damage produced by fusio neutrons is to find the „approppriate damage parameters which correlate fusion neutrons with other particles. Several damage correlation parameters such as damage energy, displacement per atom and fluence are frequently used in the literature'6,7'. However, no one of these parameters can be generalized to cover any range of irradiation for any material'7'. Recent study'8' showed that the subcascade density forming the displacement cascades gave better correlation between fission, 14 MeV neutrons and de-Be sources. This is attributed to the fact that the primary damage is the basis of radiation damage produced in any material whatever is the irradiation environment. The structural materials for the first wall of fusion reactors should meet the operation severe conditions as well as offering satisfactory and reliable properties at temperatures in the range 250-500°C which would be induced in the first wall^l. Based on the properties of the materials used in breeder reactors, austenitic 316 L type and martensitic steels have been selected as potential structural materials^10'. The martensitic steels exhibit high embrittlement under neutron irradiation at low temperature'11' which make them unsuitable for the basic first wall blanket of NET'12'. However, they are likely to have in service properties such as good dimensional stability and resistance to thermal fatigue and He-embrittlement which make them reactor relevant. Therefore, they are forseen as structural material for tritium breeding blankets to be tested in technology phase of NET. -571- The purpose of this paper is to give an evaluation of the primary damage produced in 316 L, AMCR, MANET steels, by the 17 MeV protons using the TRIM code. The 17 MeV protons are chosen since they produce an equivalent value of displacement per atom as that expected for the Next European Torus (NET) reactor*13'. 2. Materials and Damage Simulation 2.1 Materials The 3 steels used in this study have different structures and different chemical compositions. MANET type 1.4914 (German grade) has a martensitic structure at room temperature, while the other two steels, the AMCR type 0033 (French grade) and the A1SI 316 (European reference steel) have austenitic structures. The major elements and the carbon content in these steels are given in Table (1). It is important here to mention that the austentic structure of the AMCR is mainly stabilized by the 17.3$ Mn while that of the AISI is reached by the 12.34% Ni and 1.81% Mn[9J. 2.1 Simulation of Damage The computer programs for the Transport of Ions in Matter (TRIM) is used to calculate the penetration of energetic ions into solids as well as all kinetic phenomena associated with the ion energy loss such as target damage, ionization and phonon production. Details of the TRIM can be seen in references'14,15^. 3. Results The damage introduced in the studied steels was presented using a constant number of hydrogen ions accelerated to 17 MeV and incident at zero angle on the surface. The results obtained for the 3 steels were found to have the same general features, for the sake of avoiding -572- Ittm l. (Iwrtt (mttiM, M * 1 tt—mf l« t% m •• Mttl ' 4MMr «M II. » ».» J.41 Ml ».ll «wt fl.l II. t ««.! «l.l II.* $.1 •III «M I».M • I.* I.I I.fl i.un tmttm ». —tm—Jtttiti— MNI aa* %**>U r«M M«t«' aVillM—m»Af» HI in Ml f»4mt *—ft wmtmtitmtim II» in III r*ot *•*»<<• in ill III •MMUMI wtUKM*. 1 41.4 M.« •l.» •••• Mftl !• 44* | *•» ••• faMI 1. MfaM 4 \1M H«I« «•* ««a «««« «» «M(l tmt I WWII | MM •Ml | r*««l MM«f 1* MM4C*. €•* tl.l l«.l n.4 IM ••••<» l» WXIt, M» !#.• «.M ut% r«t«t «M«i«/««t#ic. *rt——4t III •* Ml) •*• Mn*«/W««4f , ItMoU *!.« •» it 10 aMf IM*«I|M. <.< • *• • II •»« MM «> III #*> *M «•• /,—1 Mirili, MIA C |»'» 1.1 «#.« ll.l »«. I M.I II l.l H.I l« II HI 4*1 «».» «•< «# »4« .1* Ml • II •»» 444 iMlilM.. "1 ... »*.* 41.» •.« 11.4 II.» W.I M.I trlt • !«•* d -1 l».» •» 144 lit • II Ml • l< II* III Ml M« WUWl, t.l II H.I II t.l t.l I.I ll.l • 1.1 II II «»/A*« II"4 -573- repetition and reducing the number of figures only those for the AISI type 316 steel are given as well as those which show significant differences. Figures (1-3) show the different paths of the ions (protons) as well as the tracks for the recoiled atoms at different target depths for the 3 steels. The damage is mostly localized in the forward direction with maximum diverting angles from the surface of 12°, 22" and 43° and mean ranges of 655, 646 and 635 urn for the AMCR, MANET and AISI steels, respectively. Almost all the ions come to stop around the mean range with maximum widths of 92, 123 and 212 im, respectively, as measured from.the ion range distribution curves (see Fig. 4). The results of the collision events of protons as given by the total target displacements and total target vacancies curves, are shown in Figs. 5,6. Similarly the recoil distribution curves as well as the replacement colisions target depth distribution curves are presented in Figs. (7,8). Table (2) gives the corresponding values obtained from the area under the curve results for the three steels. The table shows that the total target displacements/ion and vacancies/ion ranged from 135 to 198 displacement/ion and 121 to 137 vacancies/ion, respectively. The total recoils produced from the ions as well as from the primary knocked on atoms (PKA) ranged from 180 to 243 recoils. Replacement collisions amounted to 33 - 44% of the total recoils. The energy loss curves or the recoil energies curves which include the total energy to recoil and the ion energy to recoil are shown in Figs. (9,10). Table (3) gives the values of the total energy to recoils and the ion energy to recoil as obtained from the area under the curve results as well as the corresponding values of the total energy/recoil and the total ion energy/recoil for the three steels. The total energies to recoil and the ion energy to recoils were found to range from 16.7-29.6 keV and 8.28 to 10.6 KeV, respectively. The -574- TRIM-86 • Cfliwici:«n c-:"*i?n H mu ten fcw «y -• tf HeU FIG. 1 fN!M»il. i t Hc*>Ret;e! «JSu» ::c. « . •ism. • . TEIM-86 f iyiit H * m« T-2»COt. ' ' •.infers •" FIG. 2 teatcctttrt 0 v«i:« » ur :cMto>i '•••''.""• TRIM-86 H mv :«« tr.v ji ; tttNV FIG. 3 !'!«.-r.:tf: « '••g.ts.::. • « VCt' :» = IM :<.*.::.! ooo .':t t . FIGS, t-3: PATIB OF PROTONS AND TRACKS OF RECOILED ATCHS nm TIE AtCR . MANET AND AISI STEELS. -575- TKIM-86 r^^rtiION RftNC E .... H tf.V I- KMbtrfv : ttu «: FIC. 4 Trcastlt. ! 0 « t 11 .e« Vct/!«ft ' 130 g ?»:«..• •_••• • « :«::•• Tcr-gV.. : Ocv FIC. «> 0 Tretaei. s 0 Strewn} * M.0 «• •.•colon ' 130 :«UMOO.O Vs: : . ;«»•«—,1;-?rg.E* * M TRIM-86 f COLLISION EVEN' Ion: Hj( ^ Totol Target Vaconties l«n lower : 201 ionCMrfj : « FICS. 4-6: DISTRIBUTION CURVES FOR ION RANGE, TOTAL TAHGET DISPLACEMENTS AND TOTAL TARGET VACANCIES FOR TIE AISI STEEL. -576- corresponding values per recoil are 93-140 eV/recoil and 42-46 eV/recoils. The threshold displacement energy was found to be in the order of 25 eV/atom. The general features of the distribution curves are the same. The distribution curves show spectra of maxima as one scan the target depth axis from the surface of the material, Figs. (10-12). However, the positions as well as the area under the peaks vary from one steel to the other with maximum values around the mean l-ange. Table (4) gives peak positions as well as their corresponding amplitudes. 4. Discussion The results of Figs. (1-3) for the paths of protons and tracks of the recoiled atoms indicate that the damage is localized in the forward direction and depends on the alloying elements in the steels as seen from the vlaues of the mean range. However, at the endo of the paths i.e. as protons lose most of their energy, it was observed that scattering in all directions is possible up to 9 = ISO4. The average scattering angle of collision can be calculated from the theory of elastic collision since the average vlaue of COS 8 = 2/3A, where A is the ratio of the atomic number of the target material to that of photon. Using the vlaue for iron [A = 59], 9 should be around 89°. For successive collisions the average values for 8 were found to be as high as 22', 12" and 43° for the steels MANET, AMCR and AIS1, respectively and indicate that the alloy AIS1 has high scattering ability and thus the damage is not localized in a small volume if compared with that of the alloy AMCR. This is supported by the widths of the damage at the end of the paths as the corresponding values are 123, 92 and 212 pm, respectively, Fig. (4). Addition of Ni rather than Mn to stabilize the austenitic structure is therefore, responsable for the maximum scattering of the protons. Moreover, the stopping power of Ni is 14.8% higher than that of Fe and aobut 23.7% higher than that of Mn and Cr. This indicates that the alloy -577- * ««,*»**•, t* t t t * « TRIM-8'6 RECOIL DISTRIBUTION H M-.V — Tote: ncsa sutruvll«n «5t l*Kl'f" XV. tM> ft * * «" tKHT*. Jir«j;:in'{ ' M.o v* WO* {-i' '.'MZ'W.0 liAW :wkx;str : tot MO- FIG. 8 •ussa!!. : 4 JO*" vsi. :x< no = .tf...^: ikiLiy-A^ ».__.. THKE.TK'.T.H. il TRM-86 ( RECOIL ENERGIES . 0-.- n MM — TMat tnt-fi Jo Msttts . oltt .0? FIC. 9 ?*:is::::«~ 0 ~»«. «.£>« ' «SS WW •;z:.'m. TAKET OCf TH tM FIGS. 7-9: OISTMUUTIOM CURVES FOll TOTAL FUXOILS, REPLACEMENT COLLISIOHS AND TOTAL QIETBY TO IIECOILS FOR TIC AISI STEEL. -57K- .. i -V- RECOIL EKERCIES » tv M Cfr-ic-, : OtV FIG. 10 0 n» { l t » ». I.1, ,VU P.ECOIL EHERCI ES OJJ, i« 0J£ FIG. 11 o o SiSlS tt J <*• f.! g* vst; Se ::: K-.a ••Wl.O Wi inx. 01 . TI r r * ft » vrr.-i-w-s; TRIM-86f RECOIL ENERCIE5 Jo?-: H mv — ion Cfitryj io Mulls ~OP:SC:, no. 12 TAKCT Mfffl FIGS. 10-12: DISTRIDUTIOH CUI1VES FOIl 10H HOCIES TO RECOILS FOR TIIK AISI. AfCR AND HAHCT STEELS. -579- AISI should give shorter mean range which is clear from the TRIM code calculations (635 urn). The curves for collision events, recoil distribution, recoil energies Fig. (5-10) for the AISI steel as well as those for the AMCR and MANET alloys Fig. (11,12) are, in general, have similar shapes and are characterized with several peaks at definite target depths Table (4). The peaks were found to occur from the surface of the material up to the longest path of protons. This imlies that maximum effects on slowing down protons occur at target depth corresponding to certain proton energies. Thes peaks are resulting from the resonance phenomon at which the collision cross sections of protons with the target atoms have maximum values. Therefore, for the same alloy the resulting curves Figs. (5-10) should have the same peak positions and peak numbers since maximum number of displacements result in maximum number of recoils which have maximum number of energies. The resonance peaks for the three steels Fig. (10-12) and Table (4) have mixed peaks for the different alloying elements in each one. Comparison with results made on pure metals indicates that iron is characterized with the peak at 364 um. Chromium is responsable for the first peaks occurring at 608 and 624 |im. Nickel is responsable for the 225, 242 and 527 um peaks while Mn is well defined by the 15.5 um one. Overlapping peaks are located allover the path depth with maximum amplitudes and widths for those around the mean range. The results for the total target displacement/ion, total target vacancies and total recoils Table (2) showed that the AMCR steel gives lower values followed by the AISI and MANET steels. This indicated that less damage is introduced in the structure of the AMCR. However, in all cases, one would expect an average of 350 collisions to slow down protons from 17 MeV to 25 eV (the threshold displacement energy) as calculated from the average logarithmic energy decrement'14'. The differences between -580- the calculated number of collisions and observed displacements/ion (136-198) can be attributed to the fact that collisions cause displacement of atoms as well as non-displacement ones in which protons energy is responsable for rising the temperature of the first wall up to 500°Cl2l Table (2) shows that replacement collisions amounted to 43.6%, 32.8% and 42.9% for the MANET, AMCR and AISI steels, respectively. One can, therefore, expect that the total number of displacements or total number of recoils would produce equal number of vacancies. However, the number of vacancies formed are less than those of recoils indicating that part of the vacnacies were replaced by any of the atoms in the steel. Thus, part of the recoils are replaced and the rest will occupy interstitial positions. The total energy to recoil curves as well as those for ion energy to recoils Figs. (9,10) and Table (3) give lower values for the AMCR steel (16.7 and 8.28 KeV, respectively). The total energy per recoil is also low for the AMCR (93 eV/recoil) while the values for the AISI and MANET steels are 140 and 120 eV/recoil, respectively. The ion energy per recoil is nearly similar and vary from 42-43.6 eV/recoil). These values as added to the average threshold displacement energy of 25 eV give the total energy to create Frakel pairs. The energy to create Frankel pairs are 68.5, 71 and 67 eV for the MANET, AMCR and AISI steels, respectively. Diercks151 reported a value of 69.8 eV for the creation of Frankel pairs in a-iron metal. 5. Conclusions 1. The damage introduced by the 17 MeV protons in the MANET (martensitic) and AMCR and AISI (austenitic) steels using TRIM code was evaluated mainly as primary damage. The 3 steels presented respectively, 137, 121 and 130 Frankel pairs/ion, 243, 180 and 210 recoils as well as 43.6, 32.8 and 42.9% replacement collisions. -581- 2. The total energy to recoils as well as the ion energy to recoils were calculated to be 29.2, 16.7, 29.6 KeV and 6.6, 8.28, 8.75 KeV for the MANET, AMCR and AISI steels, respectively. 3. The total energy per recoil and the ion energy per recoil were found to vary from 93-140 eV/recoil and 42-46 eV/recoil. The energy to create Frankel pairs was found to vary from 67 to 71/eV. 4. The damage was found to be scattered all over the path depth of the protons depending on the stopping power and the resonance characteristics of the alloying elements in Fe (the base metal of steel). Maximum effects were found to be located around the mean range of the alloys. 5. Non-displacement collisions are strongly behined the rise of temperature of the first wall of fusion reactors due to thermal or photon energies acquired by the target atoms. 6. The damage introduced in the AISI austenitic steel is not localized and is scattered at wider angles if compared with that produced from the MANET or AMCR steels. This may be attributed to the presence of Ni atoms which have the highest stopping power compared with other elements in the steels. Moreover, the magnitude of the damage as well as its width are much smaller in the AMCR than that of the other steels due to its higher Mn content. It is recommended to use austenitic steels for the first wall of fusion reactors. References 1. M. Chazalon et al., fusion Technology, 14, 188, 82. 2. J. L. Boutard and J. Nihoul, the NET team, private coinunication, Instit. Fuer Plasm., Muenchen, Germany. 3. M. Kiritani, J. Nucl. Mat., 155-157, 1988, 113120. 4. T. Muroga and S. Ishino, J. Nucl. Mat., 117, 1983, 3645 5. R. Diercky, J. Nucl. Mat., 144,1987, 214-227. 6. H.L. Heinisch et. al, J. Nucl. Mat., 141-143, 1986, 807-815. 7. H.L. Heinisch, J. Nucl. Mat. 155-157, 1988, 121129. S. C.A. English, J. Nucl. Mat., 133-134, 1985, 71. -582- 9. G. Casini and P. Finichi, EUR 12329 EN, 1989. 10. G. Vieder, W. Daenner and B. Haferkamp, Proc. Thinteenth Symposium on. Fusion Technologv, Veresc (Italy), Sept. 1984, 2, 1363. 11. K. Ehrlich and K. Andrcko, Proc. of hit. Symp. on Fast Breeder Reactors, Lyon (France) July 1985, 231. 12. D.R. Harries, J.M. Dupouv and C.H. Wu, J. Nucl. Mater. 25, 1985, 133. 13. D.G. Rickerby and P. Finici, J. Nucl. Materials. 103-104. 1981,1577-15S2. 14. J.F. Ziegler et al, ISBN-0.08021603X, pergamann press. New York, 1985. 15. A.I. El-Shanshoury et al., 3rd Int. Conf., Al-Azhar Univ., Cairo, 6, Dec. 1993, 37. 16. I. Kaplin, Nuclear Physics, 2nd edition, Addision Wesely Pub. Co., Reading Mass, 1969,567-572. -5X3- First International Spring School & Symposium in Advances in Material Science (SAMS 94) Cairo, 15-20 March 1994 EVALUATION OF THE DAMAGE INTRODUCED IN THE MATERIALS PROPOSED FOR THE FIRST WALL OF FUSION REACTORS USING TRIM CODE PART II: ALPHA-PARTICLES EFFECTS I.A El-Shanshoury, G.I. Orabi, A.I. El-Shanshoury Department of Material* Engineering, NCNSRC, ATOMIC ENERGY AUTHORITY, Cairo, Egypt. Abstract The damage introduced by a-particles in the first wall of fusion reactors was evaluated for three different steels 316SS, AMCR, and MENT types, using the TRIM computer program. The study included a-particles resulted from the fusion reaction with energies of 3.5 MeV which escape the plasma magnetic field as well as those produced from the transmatation of B10 and Ni58. The results pointed out that most of the damage for the fusion particles is concentrated in layers with thickness ranging from 0.8 to 2.0 um with their peak effects at a mean range of about 6.1 um from the surface of the material, a-particles from the transmutation reaction of fusion neutrons, during their slowing down, with boron have a maximum effect at about 650 um from the surface of the material. These nuclear reactions can lead to the formation of helium filled cavities which have a profound influence on the mechanical properties of the first well and cause dimensional instabilities by "swelling" at higher temperatures. The results of the energy loss for the 3.5 MeV particles indicated that about 98.5% of the total energy is given to the ionization of the target materials. The balance (50-54 KeV) -585- is distributed between recoils (26-39 KeV), Frenkel pairs production (19-20 KeV) and phonon production (4.6-5.4 MeV). The recoil distribution curves and the collision events curves gave values of 267-286 vacancies/ion, and 350-400 recoils as well as 70-133 replacement collisions. It can be concluded that the damage produced by a-particles of 3.5 MeV can not be excluded relative to that produced by 14 MeV neutrons and should be taken into consideration in such microstructural evaluation for the first wall of fusion reactor. 1. Introduction Helium in the firstwall of fusion reactor is produced by two sources. The first is coming from the a-particles generated from the thermonuclear reaction which produces a-particles of 3.5 MeV energy. The second source is the reaction of neutrons of 14 MeV with the constituents of material especially, B and Nr ' ' during the slowing down process. For example, in stainless stell 316, the amount of helium generated by a fusion reactor environment is about 150ppm per megawatt years per m213'. It was reported that the enhancement factor for helium generation in a fusion reactor compared to a fast breeder reactor is about 30 at the same level of displacement damage'31. One of the most severe effects of helium is the degradation of the mechanical properties of material at high temperature, which at present would probably limit the operating temperature of fusion reactor'41. The primary mechanisms responsible for that are generally associated with helium and self-point defect clusters such as helium vacany and helium interstitial clusters, helium stablized materix cavities, helium assisted interstitial loops and helium-stabilized grain boundary cavities. Except for the last type, the clusters listed above are generally believed to produce matrix hardening rather than grain boundary embrittlement'5'. This depends on the temperature at which -586- the material is operated and the initial state of material which is characterized by its, chemical composition and microstructural parameters such as grain size, dislocation density and precipitate density. Limited studies*5"7' showed that helium in conjunction with displacement damage may drastically reduce the ductility properties, change the failure mode, and also alter the microstructure of the alloys of interest in fusion reactor technology. It was found that the extent of this degradation is related to the He/dpa ratios, where dpa represents the level of damage produced during fusion irradiation. In order to increase our knoweldge of microstructural and chemical composition effects on the structural damage produced in the first wall of fusion by the a-particles of 3.5 MeV, a computational study using TRIM code* ' is done on three different steels AISI 316, AMCR, and MANET which were selected for the manufacturing of the first wall of the future fusion reactors. The 316SS (Fe/Cr/Ni) and AMCR (Fe/Cr/Mn) are austenitic stainles steel while the MANET (Fe/Cr) is a martenstic steel. 2. Materials and Damage Simulation Three different materials 316SS, AMCR and MANET have been used in this study. Their chemical composition is listed in Table 1. The main difference exists in Ni and Mn contents. The damage produced in these steels by 3.5 MeV a-particles is stimulated by utilising the computer program for the transport of Ions in Matter (TRIM). The content of this program and its facilities are published in18,91. 3. Results The damage introduced in the studied steels was presented using a constant number of a-particles (Helium ions) acelerated to 3.5 MeV and incident at zero angle on the surface. The results obtained for the 3 steels were found to -587- have the same general features, therefore, for the sake of avoiding repetition and reduction of the number of figures only those for the AISI type 316 steel are given as well as those which show significant differences. Figures (1-3) show the different paths of the particles as well as the tracks for the recoiled atoms at different target depths for the 3 steels, The damage is localized in the forward direction with maximum diverting angles from the surface of 30°, 15° and 12" for the AISI, AMCR and MANET steels, respectively. However, most of the damage is confined to angles of 10°, 17° and 12°, respectively. The corresponding mean ranges for the 3 steels are 6.03, 6.11 and 6.03 pm, Almost all particles come to stop around the mean range with maximum width of 2.01, 1.20 and 0.85 mm, respectively, as measured from the ion range distribution cruves (see Fig. 4). The results of the collision events., recoil distributions and replacement collisions of the particles are shown in Figs. 5-8. Table 2 summarizes the corresponding values obtained from the area under the curve results for the 3 steels. The energy loss curves for the decelerated a-particles as well as the recoiled energetic atoms are given in Fig. (9-12). Fig. 9 shows the ionization-target depth curve which is typical of the attenuation of a-particles in matter. Intense ionization begins from the surface of the target material and slowly increases reaching a maximum and then suddenly decreases as the particles come to stop. The area under the ionization depth curve gives the total energy lost by the a-particles. Generally 98.5% of the total energy of the particle is lost in ionization. The rest of the energy is distributed between recoils, vacancies and phonons. Table (3) gives the corresponding values for the 3 alloys studied. It gives the values concerning the total energy to recoils, the ion energy to recoils, the total energy to vacancies, the total energy/recoil and the ion energy/recoil. It includes also the total energy/vacancy as well as the energy to create Frenkel -588- Ttbtl I. Chtnlfl Competition. M t C<>.4t«<> r« c* u/ Us tin e Slttt «wer ir.s r«. l ».» CO . Ml *. if HMCI n.« i«. i <». r <«. 1 ir. j «.i MSI »« ir.«« II. J« I.J i. it t.esit iAfi£ 7. WIJU-S rai CUXISICM aims «n Htcon. oismmniuis SIEEI. MISt men WWET Total target dlsilaeonatts/tai 290 5» 2GI Total Target vacaiiclcs/lon 2S6 267 286 Total recoils 556 100 380 Replaccncnt Collisions, X 20 55 • Mean Range, ya 6.05 6.11 27 6.05 TADLE 5. VALUES IF IOTA. ENERGY AND ION ENERGr TO RECOILS STEEL AISI AIKR IWHET Total energy to recoils, KeV 39.2 11.2 53.1 len energy to recoils, Kcv 13.95 15.5 11.3 Total OKryy to vac'ii.lcs, KRV /.25 7.5 7.1 Total cucrgy/rccoll.cV/rccoll 110' 105 III) lot aiergy/rccoll, ev/rccoll 39.2 SJ.75 37.G Total crcrgy^maicy cV/Vac. 25.5 PB 21.8 Bieroy to create Frafccl Pair, cV CI.5 06.75 02.1 -589- — ,'*.<.'.•;::. : ;r. tett<z* -• "•:". - .'. — v35«lCttor. Cimc-tTt — -jri^i3li«5crasu:t- AIJI % l^ 10 um TAWET DiFTH V) -M FIG. 1 Fig. Z Fig."3 FIG. 1-3: PATHS OF 3.5 McV o<-PARTICLES AND TRACKS OF RECOILS ATOMS IN AISI 316, AMCR AND MANET STEELS TRIM-86 AT* •<"*••*'• •vr--- ION' RANCE ivJJ.HC n» ;.e.tn To.-ae: •t. 1 J.. Law-. • IWftMT* * m lof. ttieft t J« PICA ION RANGE DISTRIBUTION frttf 0 CURVES IN AISI 316 STI-HL treur.::. i » tWwk Jtf :*-IHJ' •: :.. vc: s J »!*.•. 1 . 1 r/AsrK'w :; 4-:-•:.--atu : u t! »: ?'jtf ^sasj Nlj IBUSS^ i RECOIL DISTRIBUTION "1 |RE?UCEKEIIT COi-LIi»0«!5 • r ?! / ll II t I % . !l '-•••-••"•• \„ vAM -e. T»KCT M«*« U"irrbi ,1* TAMP rem 10*. THUnKTTN ttt* JIG. S FIC. 6 FIG.7' FIG. t FIC. S-8: COLLISION EVI-NTS. RECOIL DISTRIBUTIONS AMI RF.n.ACKKP.ST COLLISION'S OF 3.i M..V £K .PARTICLES IN AISI JM.SS TADI.n 4.' Target Depths at Peaks Fur I lie irradiated Slccls a* Obtained Proa the Recoil Entrlf Distributions (a) AIS1 Pt»k Depth, O.SS 3. 71 <..i>S S.47 Ifml Amplitude , 2.71 !•!« 1 .77 •22 S.91 | eV/A i (Ii) AMCR Peak Depth, 1.56 .97 4.35 5.23 S.47 6. OS (Jim) ABflitude, 0.83 1.27 O.B« 0.9) j 1.55. 6.58 eV/A J__ • (c) MANET Peak Depth, 3.2 3.6.' 4.49 5.08 S.4 COS (/m) Amplitude 0.98 0.6T 0.13 o.ea 1.45 6.2S eV/A IADLF 5. EHHRGY LOSS. X AISI AHCR HA1IE! Ionization 9S.55 W.'i? •B.57 Fhonons 0.13 0.15 0.1S Recoils 1.12 1.13 1.U9 Vacancies 0.20 0.22 0.19 •592- •;'XM>.V.W •»!»•(! -.rwn I ». » : » IONIZATION /ION RECOIL ENERGIES 'RECOIL 'ENERGIES FH0N0NS / ICN •iT.' . - *4 J. v IS. .li .1 v© •A |rS:i-r::'.-f' ;i >iv"-'--** J il 1 jJ &= T*-:ET st»ni FIG. 10 FIG. 12 FIG. 9-12: IONIZATION, ENERGY LOSS OF 3.S MeY C< -PARTICLES AND RECOILED ENERGETIC ATOMS CURVES FOR AISI 316 STEEL at definites depth Fig. (5). The peaks were found to occur from the surface of the wall material up to the longest part of the particles. These peaks were identified to characterize the elements of the alloys. Iron is responsible for the 2.97 urn peak as well as the 5.05 urn peak which is the additive peak for all elements. Ni is significantly known by the 0.35 urn peak while Cr is well indicated by the 5.47 mm one. Mn, however, has more than one peak occuring between 4.49 and 5.08 um range. Those peaks are the resonance ones at which the collision cross sections are maximum. Therefore, one would expect maximum displacement of atoms, maximum recoils, maximum vacancies as well as maximum values for the corresponding energies as seen from (Figs. 5-8 and 9-12). The results for the total target displacements/ion, total recoils Table (2) as well as those for the corresponding energies Table (3) do not show any significant differencies. The only considerable observation is that for the value of replacement collisions for the AMCR stell. Table (2) shows that 33% of the collisions were replaced in the structure of the AMCR steel versus 20% and 27% for the AISI and MANET steels, respectively. This could be attributed to the MN content of the alloy (17%) if compared with the other steels. Moreover, the percent of replacement collisions for the particles are in general far less than those of the 17 MeV protons given in the previous paper'10l The results of ion energy to recoils curves, energy to vacancies as well as the total recoils and vacancies are given in Tables (2 and 3). It is seen that the ion energy per recoil is around 38 eV/ recoil, but the ion energy per vacancy ranges from 25-28 eV/Vac for the 3 steels. Consequently, the energy required to create Frenkel pair ranges from 62.4 to 66.75 eV which falls in the same order of those obtained from proton results'10' and reported by Direcks for -iron'11'. The most signficant results which appeared from this study are the intense ionization of the 3 alloys under the deceleration of a-particles, Fig. (9). It is quite clear that -594- about 98.5% of the energies of the a-particles are lost in the ionization of the 3 steels as shown in Table (5). The balance of the energy is distributed as 0.15% for phonons, 0.2% for vacancies and about 1.1% for recoils. Neutrons, however, do not produce ionization and most of the energy is lost during the slowing down by the collision process. Therefore, one can conclude that the responsible particles for the rise in temperature of the first wall of fusion reactors are neutrons rather than a-particles. It is important at the end of this paper to give a general picture for the damage produced by the particles libarated from the fusion reaction. Damage by a-particles (helium gas as the particles come to rest) is mainly localized at about 6 urn from the surface of the first wall with widths between 0.8-2 um while neutrons damage is localized for away at about 650 um from the surface with widths ranging from 92-212 um. This means that two regions of damage are produced in the first wall during fusion irradiation. One is very near to the surface and is a location of embrittlement by helium bubbles while the second region in more deep, more thick and is considered as privilage site for microstructural modification. Comparison with other studies^12' which showed that during irradiation of AMCR steel by a-particles at room temperature, interstitial clusters were observed. E. Ruedl and Coworkers'13^ found further that these clusters were enriched in chromium up to 40%. Apparently Cr diffuses faster during irradiation by an interstitial diffusion mechanism than Fe or Mn. This behaviour is looking different from that found in nickel containing steels in which nickel containing steels in which nickel diffuses faster than Fe or Cr during irradiation with high energy particles. Consequently, the matrix in Fe-Cr-Mn steel becomes depleted in Cr in contrast to Fe-Cr-Ni alloy in which the matrix becomes depleted in Ni'12'. Finally, it is important to mention that helium produced by transmutation reactions -595- Pairs. Table (4) shows the distribution of the energy loss (as percent) over the four significant processes of ionization, phonons, recoils and vacancies. In general features of the distribution curves are the same. The curves show spectra of maxima as one scans the target depth axis from the surface of the material, Figs. (5-8) and (10-12). However, the positions as well as the area under the peaks vary from steel to the other with maximum values around the mean range. Table (5) gives the peak depths from the surface of the materials as well as their corresponding amplitudes. 4. Discussion The results of Figs. (1-3) for the paths of a-particles and tracks of the recoiled atoms as well as the diverting angles of the particles indicated the localization of the damage in the forward direction for the 3 steels. Moreover, the values of the maximum angle of diversion show that AISI 316 and AMCR steels have higher scattering power than that of MANET steel (30% 15 and 12°, respectively). This is supported by the widths of the damage at the end of the paths since the corresponding values are 2.0, 1.2 and 0.85 um, respectively; Fig. (4). Addition of Ni in the AISI rather than Mn in the AMCR to stabilize the austenitic structure is therefore, responsible for the maximum scattering of a-particles. This is quite clear from the significantly recorded peak at 0.35 Jim from the surface of the AISI steel Fig. (5-8). Moreover, although the stopping power of Ni is higher than the other elements (Fe, Mn, Cr Mo), its effect is not observed in the mean range values of the a-particles in the 3 steels as they fall around 6.03 |.im. Proton results'10', however, is more sensitive than those given by a-particles. The curves for collision events, recoil distributions, recoil and phonon energies (Figs. 5-8 and 9-12) for the AISI steel as well as those for the AMCR and MANET steels have generally the same shapes and are characterized with peaks -596- with n-14 MeV is not taken into consideration in this analysis. This helium is principally related to the chemical composition of steel in particular Ni and B elements. The fcrritic steel like MANET has very little nickel relative to AISI and AMCR steels. So low level helium is expected in MANET steel. The impact of this helium on the performance of 316ss, AMCR & MANET steels as first wall for the future fusion reactor is not therefore the same. 5. Conclusions Following this computational study, the impact of a-particles of 3.5 MeV on the state of the 3 steels 316ss, AMCR and MANET proposed for the first wall of future fusion reactor can be as follows : 1. Most of the damage in all steels will be close to the surface of the first wall with thickness ranging from 0.8 to 2 um. Their peak effects situated at a mean range of 6.1 p.m. In the case of 17 MeV protons the damage was concentrated in the bulk of the first wall at about 650 um from the surface. 2. The steel type 316ss presented, relative to the AMCR & MANET steels, the largest scattering phenomenon. This is owing to the higher Ni content of 316ss. 3. The AMCR steel which contains higher Mn content presented the largest replacement collision during the a-bombardement against 20% for AISI 316 and 27% for MANET steel. 4. The majority of the energy of 3.5 MeV a-particles (98.5%) is lost in the ionization process for all the 3 steels. This is the contrary to what happened in the case of neutrons irradiation where most of the bombardment energy is losted in the collisuon process. 5. The damage produced by fusion environment can be distributed in the first wall of fusion reactor into two regions. One is very thin "(<2|im) and near to the surface and is considered as privilage site for helium bubble formation. The second region is thicker (92-212um) and -597- situated far away from the surface and is a location of severe structural damage. 5. References /. H. Trinhaus, Radiation Effects, 101,1986, 91-107. 2. P. Jung, J. Nucl. Material, 144,1984,43-50. S. H. Schrotder & P.Batfalsky, ASTM STP 870, 1985, 745756. 4. D.R. Harries, J. Nucl. Material, 82,1979.2-10. 5. L.K. Mansur & M.L. Grossbeck, J. Nucl. Mat. 155-157, 1988, 130-147. 6. P.J. Maziazi, ASTM STP 979,1988, 116. 7. G. Orabi, P. Finici, Technical Note, ISPRA, 1989. 8. 1. El-Shanshoury et al., 3rd Int. Conf. Al-Azhar Univ., Cairo, Dee. 1993, 37. 9. J.F. Ziegler et al., ISBN-O. 08021603X, pergomann press, New York, 1989. 10. LA El-Shanshoury et al., SAMS 94, March 1994, Cairo, Egypt. 11. R. Dierchs,J. Nucl. Mat., 144,1987, 214-227. 12. W. Schule, EUR 12335 EN, CEC 1989, ISPRA, Italy. 13. E. Ruedlet al., Phys. Stat. Sol. (A). 107,1988, 745. -598- EGtZojSHo First International Spring School & Symposium on Advances In Materials Science ISAMS 94) 15-20 March 1994 Study ofElcctroilc Kinetics at the Interface between High-T, Superconductors and Solid Electrolytes HulaS. Zr-tlihml. M. A. AhiM KauuJ. II'. J. hinn:' Physics Department. Faculty of Science. Am Shams University. Cairo. Egypt • Institute of Physical Chemistry and Electrochcmisuy. Karlsruhe University. FRC. ABSTRACT Four different contact methods have been utilized to study the electrode kinetics of high-Tc superconductor (HTSCVsolid electrolyte (SE) electrochemical systems. By using a transient technique, cell currents were measured versus time at a constant potential. Investigations with different electrochemical syslems were implemented to calculate the potential dependent charge transfer resistance of the electrochemical phase boundary reaction at the used solid electrolyte, the contact resistances and the olimic tesislaiice of the solid electrolyte in contact with HTSC. from the dc limits of lire transients, a small increase of the electrode kinetic current is observed around Tc which must be correlated to superconducting state ol HTSC. The results are interpreted as an enhancement of the electrode kinetics according to the values of the parameters mentioned before 1. INTRODUCTION Electrochemical experiments on different interfaces, formed by condensed matter, have been studied extensively in the last few years, to study the transfer of charge carriers across different interfaces. Electrochemical studies on ceramic high-temperature superconductor (HTSC) and ionic conductor (IC) interfaces in a wide range of temperatures are of great interest for both kinds of ionic conductors. Electrochemical charge transfer process can be -599- studied with either liquid or solid electrolytes. Liquid electrolytes (LE) hav the advantage to contact the substrate perfectly, to exhibit a well studied electrochemical double layer structure. Unfortunately, the choice of inorganic or organic solvent mixtures being liquid at actual Tc's and dissolving a supporting electrolyte to maintain a sufficiently high ionic conductivity, is restricted to only a few systems [1 J. on the other hand, obvious disadvantages of solid electrolytes include contact problems with HTSC, largely unknown structure of the electrochemical double layer, and low ionic conductivity at low temperatures. However, it is easy to prepare thin electrolyte layers in order to reduce the ohmic resistance. Electrochemical experience on HTSC/LE interfaces were mainly carried out at elevated temperarures, i.e. T » Tc, in order to study surface stability, aging effects, phase boundary reactions such as corrosion, metal deposition and redox reactions, photo-electrochemical effect, etc. [2-5]. The mechanism of the charge transfer process across HTSC/LE interface can only by studied at low temperatures T « Tc. in order to get information about the electrode kinetic current and what happens to this current when the system Ag/HTSC/SE/Ag is in (he superconducting state and what is the effect of the contact resistance. This paper deals with the explanation of the above questions. Results on four different contact methods with electrodes are given. Investigations of the charge carriers transfer phenomenon in the point of view of the current electrode kinetics were presented. 2. EXPERIMENTAL SET-UP The electrochemical measurements of the electrochemical cell Ag/HTSC/SE/Ag were performed with press contact, i.e. mechanically, between polycrystalline HTSC as a working electrode and RbAgJs as a counter and reference electrode by four different types of contacts, by sputtering silver (20 nm diickness) on the backside of the SE, silver foil (0.1 -6(H)- mm in thickness), silver powder and special silver epoxy resin for low temperatures (polytec). AU external contacts of the HTSC as well as to the .inter electrode and reference electrodes were made using the silver epoxy resin as a contact material for metallic silver fibers. The methods used are the four-probe technique for the determination of the bulk resistance of the HTSC and the SE contact and interfacial resistance. In order to work in an appropriate temperature range, a cryostat based on a liquid-helium cooling system (Cryophysics LTS-22-DRC 91C) was used to adjust the temperature in the range 12K < T < 298K under vacuum conditions with an accuracy of IK. Two different techniques were applied potentiostatically at the rest potential AE. The first is the electrochemical impedance spectroscopy (EIS) technique in the frequency domain followed by a transfer function analysis. The second is a special transient technique in die time domain using a large signal system perturbation with a periodic rectangular function of the potential. Both techniques gave identical results, however, the EIS measurements used to confirm the results from transient technique. 3. RESULTS AND DISCUSSION Electrochemical measurements were performed under exact potential and temperature control to get the ohmic resistance RQ. of the superionic conductor. The values of R were found to be rather high at low temperatures. However, they were low enough for electrochemical studies on the interfacial behavior of HTSC/SE cell at temperatures above and below Tc. Accordingly, the increase in the contact resistance. Rc, increases the electrochemical measurements Z(s), so, utilizing different contact methods gives the possibility to select one of the optimum results. The temperature dependence of the four contact resistance is shown in Fig. 1. It is seen form -601- thy figure that »h?. contact resistance Rc(Ag/SE) is sharply increased with •Jr: "easing temperature. So, Rc becomes higher and higher at low •.crtiperuiures, winch is the temperature range pose. Also, results show that the o.i'-z-ici using silver sputtering and silver foil are of lower contact resistance 'Jim using silver powder and silver paint. Indeed, this difference in Rc values with different contact methods was related to the effective area of the deposited silver on the contact surface. The transfer process of the charge U rN 10 i £ u u -.— Ag sputtering J_ -•— Ag foil "CTI o _ • _ Ag powder -•% Ag paint Q,1 •••!• /2 1£ 9'6 1 120 temperature/K Fig. ! Four different contact rcsisianc-s of uic system Ag/RbAgjI/Ag. •urn.::? across the interface HTSC/SE was studied via transient technique by ..'.v.pioying an alternating potential with amplitude IV between the working and :;v.: counter or reference electrodes. The current transients were recorded versus time. The dc limits of the cathodic and anodic current transients were found to be nearly syrrtmetrical with respect- to / = 0. Alternating pulse polarization was chosen in order to avoid irreversible polarization effects within the SE. the difference of the anodic and cathodic, quasi-dc current -602- IAJI, measured after x = 20s, is logarithmically plotted versus T in Fig. 2. for the polycrystalline high-Tc superconductor Tl2Ba2CaCu2Og measured by two-probe technique. Two principal features of this result are apparent, the first is the strong decrease of IA/1 with decreasing temperature which is observed at T » Tc corresponding to the increase in the resistivity. Another important feature of the results is that around Tc, an increase in the current is observed which coincides with the transition or critical temperature of the suprconductor. The electrode kinetic current increases by 1.4 times around the transition temperature. However, die increase is not maintained, and when 3* a 4 6 8 10 — 12 14 Fig. 2 Temperature dependence of the current density measured by 2-prdbc transient technique in the jystemTUBi;CaCu;(yAg" glass. the temperature has dropped approximately another 10 K the current density resumes (he behavior it would have had if superconductivity had not set in. It is easy to give a hypothesis as to the increase of the current density at the superconducting state. Thus, here the Cooper pairs which easily perform transitions through'the crystal come to the interface, since the coherence length of Cooper pairs in HTSC's is about 0.6 + 0.1 nm, and their properties of minimizing interaction of the surroundings are maintained so that the ease of electron transfer across the double layer is increased. -603- However, it would have been expected that this enhanced current would have continued as the temperature fell down, but this is not the case. Similar results have been observed by Bockris and Wass [6] for the hydrogen evolution reaction in the system YBaXujO^frozen HCI04.5.5H20. Pinkowski et al. [7] observed hump of the exchange current density in some redox reactions examined in silver 3"-alumina. Murray [8] also reported a non-monotonous temperature dependence of the double layer capacitance around Tc in the system Tl2Ba2Ca2Cu2O,0/mixture of liquid organic electrolyte with low melting temperature. Therefore, the extraordinary exchange in the current density would correspond to either decrease in the charge transfer resistance, R, (kinetic effect) or decrease in the ohmic resistance Rn (proximity effect) [9J. To distinct between the above two proposals, the electrochemical experiments of the system Ag/HTSC/Ag, Ag/SE/Ag and Ag/HTSC/SE/Ag are considered. The experimental overall impedance of the later system consists of a series combination of the following components: Z = ZHTSC + ZAg/HTSC + ZSE + ZsE/Ag + ZftTSC/SE (1) The first four components were determined using 2- and 4-probe techniques of the first two systems as shown in Figs. 3 and 4. The results show a relation between the resistances of HTSC and SE, respectively, as a function of temperature. In both figures, curve (a) gives the bulk resistance of the material Rfi using 4-probe measurements : lim{Z(s)} = Rn(HTSCorSE) (2) Curve (b) represents the 2-probe measurements which include the contact resistance term given by : -604- '80 85 90 95 100 105 lid 115 Temperature/K Fig. 3 Temperature dependence of RbAgJj impedances, bulk resistance (a) measured by 4-probe technique, overall impedance (b) measured by 2-probc technique and conucr resistance (c) calculated from (a) and (b). Fig. 4 Temperature dependence of the HTSC resistance, bulk resistance (a) measured by 4-probe technique, overall impedance (b) measured by 2-f robe lechnique and contact resistance (c) calculated from (ai and (b>. -60S- lim{Z(s)> = 2Rc (Ag/HTSC or SE) + Ra (HTSC or SE) (3) Curve (c) was calculated from (a) and (b) to get the contact resistance of both HTSC and SE with silver. The dc currents of mmsQ and IAzlAg axe given at a constant potential AE between the working and reference electrodes approximately by: (A/lj=AE/(Ri+R{ + RU (4) J j Where. RC . R, and Rn\ denote the contact resistance between j= HTSC or Ag and the solid electrolyte SE, die charge transfer resistance at the interface and the ohmic resistance of SE in contact with j, respectively. According to the first proposal, mentioned before, Rt decreases sharply at HTSC mSC T < Tc. IA/I in equation (4) is determined by the values of Rn directly only under the condition : RHTSC< which is observed in the electrochemical impedance measurements shown in HTSC Fig. 5. On the other hand, the second interpretation presume that Rn drops down at T < Tc due to proximity effect. This leads to an increase of HTSC lAil under the condition : Rgrsc«R{1TSC-Rinrsc. (6) which was never seen in our measurements. Therefore, the extraordinary exchange in the current density is interpreted as a real quantum electrochemical kinetic phenomenon [9] caused by the transfer of the charge carriers across the interface HTSC/SE. -606- T/K Fig. 5 Electrochemical impedances measurements vs T on the system Tl:Ba:CaCu:0,/RbA£4l,. 4. CONCLUSION The present investigations emphasize that the change in the number of charge carriers transferring across the interface HTSC/SE, around Tc affect the electrode kinetic causing an extraordinary enhancement in the current density. Using different contacts to confirm that the non-monotonous behavior of the current density was interpreted as a real quantum electrochemical kinetic phenomenon. -607- REFERENCES [l] A. Pinkowski, J. Doneit, K. Jiiitner, VV. J. Lorcrtz, G. Saemann-Ischenko. T. Zettcrer and M. W. Breiter, Electrochemical Acta 34 (1989) 1113. [2| M. Bachtler. W. J. Lorenz, W. Schindlcr and G. Saemann-Ischenko. Modem Phys. Lett B2 (1988) 819. [3] M. Hampel, E. W. Grabner. M. Bachtler and W. J. Lorenz, Modern Phys. Lett B3 (1989) 303. (4J J. M. Rosamilia and B. Miller. J. Electroanal. Chem. 249 (1988) 205. 15] S. Rochani. D. B. Hibbcrt, S. X. Dou. A. J. Bourdillon. H. K. Liu. J. P. Zhou and C. 0. Sorell. J. Electroanal. Chem. 248 (1988) 461. [6) J. O'M. Bockris and J. Wass. J. Electroanal. Chem. .267 (1990) 329. (7] A. Pinkowski, J. Doneit, K. Jiittncr, W. J. Lorenz, G. Saemann-Ischenko and M. W. Breiter, Europhys. Lett. 9 (1989) 269. [8] R. W. Murray, unpublished lectures at the University of Karlsruhe (1990). (9) H. S. Zaghloul, M. H. Zayan and M. A. Abdel-Raouf. Proceedings of the International Symposium on High-Tc Superconductivity and Its Applications. Cairo, Egypt (1993). -WIS- EG f£o /I£M First International. Spring School & Symposium on Advances in Materials Science (SAMS 94) 15-20 March 1994 ADVANCED AEROGEL MATERIALS FOR PHOTOCATALYTIC DETOXIFICATION OF CYANIDE WASTES IN WATER Mohamed Sayed Ahmed Department of Metallurgical Engineering, College of Petroleum and Mining Engineering, Suez, Egypt ABSTRACT Ultraviolet (UV)-irradiated Ti02-Si02 aerogels were used in this study as catalysts for oxidation of cyanide species (CN") in water to C02 and N2. Ti02-SiO2 aerogel catalysts were prepared by the sol-gel and supercritical drying techniques. Three types of this aerogel were prepared with different Si02 content (i.e., Ti02:SiO2 molar ratio = 1:1.3, 1:2.6 and 1:3.9). It was observed that with increasing Si02 content, shrinkage and apparent density decreased, however, translucency increased. This resulted in higher photocatalytic activity for oxidation of CN" in dilute solutions (1000 and 522 ppm) of ferric cyanide. Compared : with TiOj powder (i.e., anatase with Ti02> 98% and particle diameter < 1 urn), it was proven that the aerogel with higher Si02 content has higher efficiency in the photocatalytic oxidation of CN" at similar experimental conditions. This aerogel catalyst was completely recovered and reused several times to check the activity loss. No significant change in the photocatalytic activity was observed. 1. INTRODUCTION Hazardous industrial aqueous wastes containing free and complexed cyanides are generated in lai^e volumes annually in metal finishing industries including electroplating and heat treating operations.1'2 Unit processes typically used frr treatment of these cyanide wastes are: alkaline -609- chlorination, electrolytic oxidation, ozonation and to lesser extent, wet air oxidation and ion exchange. Currently, alkaline chlorination is the best available proven technique for treatment process.2'3 However, this cyanide waste treatment technique has two important limitations. First, only the free cyanide is oxidized leaving behind the complex metal cyanides in sludge. The sludge thus formed has to be disposed of as hazardous solid waste, usually in a landfill. Second, a highly toxic (as toxic as hydrogen cyanide) cyanogen chloride gas can be formed and released during the treatment process. Recent investigations showed that semiconductor powders, such as Ti02 (anatase), suspended in solution can utilize solar energy for photo- catalytic oxidation of cyanide species (CN") .3"6 Irradiation of a Ti02 particle with UV-light of wave lengths of 300 to 400 nm, from either a xenon light or unfocused sunlight, causes excitation of an electron (e) to the conduction band leaving a hole (h+) in the valence band. Apparently these electrons and holes can trapped in surface levels with recombination being sufficiently slow that the electron transfer process can occur. The conduction band electrons reduce oxygen while the valence band holes cause the oxidation of CN". Thus the process of removal of CN" can be represented as follows: + Ti02 + (hv) = Ti02(h + e) surface 0.5O2 + 2e + H20 = 20H" 20H" + 2h+ = 20H- CN- + 20H- = OCN" + H20 20CN" + 02 = 2C02 + N, Overall reaction: „.Q 2CN" + 202 + H20 •- 2C02 + N2 + H20 UV-light Where h is Plank's constant and v is frequency.7 As the photocatalytic oxidation reaction takes place at interface between the solution containing CN" and Ti02 semiconductor, the rate of oxidation -610- will depend on surface area of the semiconductor. 3 Therefore, colloidal Ti02 sol was used. However, the main disadvantage of using colloidal Ti02 is the difficulty to recover all Ti02 used after the treatment process, i.e., the effluent stream will be contaminated with Ti02. Alternative methods for synthesizing colloidal Ti02 semiconductor aggregates involve the use of solid matrices such as porous glasses.8 The size and shape of the semiconductor are determined in this case by the morphology and texture of the microscopic voids of the host template where aggregation occurs. A novel concept described ir> this paper uses the sol-gel technique to prepare Ti02-Si02 aerogel (porous and translucent matrices) in which Ti02 acts as the semiconductor in photocatalytic oxidation to destroy complex metal cyanides, e.g., ferric cyanide, in aqueous solutions. Such aerogel materials can have over 90% porosity, very low density and extremely high surface area, currently several hundred m2/g and a pore diameter distribution extending from nm to um.' So, the primary objective of this work was to determine whether the Ti02-si02 aerogel samples prepared in the laboratory can effectively be used for photocatalytic oxidation of ferric cyanide in aqueous solution. Once this has been realized, research will continue to refine the aerogel preparation and to optimize the oxidation process. 2. EXPERIMENTAL PROCEDURES 2.1. preparation of the Photocatalvsts In the first two series of tests two different Ti02 powders were used. The first powder (Ti02 > 98%) had particle diameter < 1 um and the crystal structure was primarily anatase. In the second series of tests rutile (Ti02= 93.3%) with particle diameter < 200 um was used. In other three series of experiments, Ti02-Si02 aerogels with different compositional molar ratios were used. The tested aerogels were manufactured by preparing the corresponding alcogels using the sol-gel processing technology, then by super critical drying to remove the alcohol. -611- (a) Synthesis of Ti02-Si02 Alcogels: Three samples of Ti02-Si02 alcogels were prepared with different TiO2:Si02 molar ratios of 1:1.3, 1:2.6 and 1:3.9 (i.e., 1:1, 1:2 and 1:3 wt. ratio) respectively. All gels were synthesized by hydrolysis and polycondensation reactions of organometallic compounds. Tetraethyl-orthosilicate (Si(OC2H5)«) and titanium-isopropoxide(Ti(OC3H7),) were used as sources for Si02 and Ti02. Gelation took place in the presence of ethanol (C2HsOH) which is readily soluble with the organometallic compounds and the distilled water (H20) used for hydrolysis. Nitric acid (HN03) was used as a catalyst. Details of the process used to prepare the appropriate alkoxide solution of TiO2-Si02 are published elsewhere.10*11 Figure 1 shows a flowchart of the sol-gel procedure used. (b) Supercritical Drying of TiO2-Si02 Alcogels: Drying of the Ti02-Si02 gels was carried out under supercritical conditions. For ethyl alcohol, supercritical conditions are 243 C and 63.6 bar.10 For this reason the gels were introduced into ?n autoclave with excess amount of alcohol to attain a pressure higher than the critical pressure at the critical temperature. After tightly closing, the autoclave was flushed twice with 2 bar dry N2 and finally pre-pressurized up to about 10 bar N2. The temperature of the autoclave was increased, at a rate of 1.0-1.5 C/min, to a temperature of about 280 C and the pressure in the closed system rose to about 110 bar. After stabilizing the maximum temperature and pressure for about 3 hr, vapor outlet micrometering valve of the autoclave was slowly opened to vent out the solvent to a condenser. All the solvent from the autoclave was evacuated in 3-4 hr at a constant temperature of about 280 C. After reaching the atmospheric pressure, in order to remove the trapped solvent vapor molecules, heating was continued for about 20 min and finally the autoclave was flushed twice with 2 bar N2. During this treatment, condensation of the outlet solvent vapor on the walls (pyrex glass) of the condenser was clearly seen. Then, the aerogels were heat treated in air at about 550 C for about 5 hr. The aerogel monoliths were then gently crushed to granules of grain size < 3 mm. -612- Si(OC2H5)4 + C2HsOH H20 HNOj (1 mol.) (4.5 mol.) (1 mol.) (Few drops) Reflux at 50-60 C during stirring. Cooling down to room temperature. Addition of Ti(OC3H7)4 during vigorous stirring. (Exothermic reaction) Cooling down to room temperature, Addition of the remaining H20 during stirring. Pouring into a mold and covered for gellation. Fig. 1. Flowchart of the sol-gel procedure used to prepare Ti02-Si02 alcogel. -613- The Ti0j-SiO2 aerogels were characterized by measuring the linear and volume shrinkage and determination of the apparent density. * Linear (diametrical) and volume shrinkage, L/L0 and V/V0 respectively, were measured for each aerogel monolith after drying by applying the following relations:" L/L0 = (L0 - L) / L0 V/V0 = (V0 - V) / V. Where L and V are the diameter and volume of the aerogel monolith after supercritical drying, and L„ and V0 are the diameter and volume of the corresponding alcogel (i.e., before supercritical drying). * The density of the aerogel could not be determined by the normal Archimedes method because of it porosity. Therefore, the density (/2erogel) was determined by applying the following relation:13 3 ^erogai = 0.997 Ma / (M0 - MJ , g/cm Where Ma= weight of the aerogel sample in air (dry weight), M0= weight of the aerogel sample after immersion in a hydrophobic oil for several hours, M„= weight of the aerogel sample in water with filled pores by oil, and 0.997 = density of water (g/cm3) at 25 C. 2.2. Photocatalytic Oxidation of CN~ Dilute solutions (1000 and 522 ppm) of ferric cyanide, one of the most common species present in industrial cyanide wastes, was used in this study as a synthetic cyanide waste. All experiments consisted of illuminating about 300 ml of the cyanide solution in a reactor (pyrex glass) containing the catalyst. The catalyst loading was maintained constant at 1.0% Ti02 by wt., in the form of Ti02 powder (i.e., anatase or rutile) or Ti02-Si02 aerogel granules. The catalyst was kept in suspension by bubbling air through the solution during the entire reaction process. The photoillumination was provided by four commercial blacklight-blue fluorescent bulbs (Ge BLB - 15 W), arranged parallel to the reactor, as shown*in Fig. 2. The illumination wavelength of this light was predominantly 320-440 ran. Open to air Solution containing Fe(CH), + Catalyst Fig. 2. The photocatalytic reactor. Control tests, with irradiated solution in the absence of catalyst or catalyst containing solution in the absence of irradiation, were carried out also. Periodically, about 10 ml sample of the reacting solution was withdrawn by syringe and analyzed for concentration of ferric cyanide, using a spectrophotometer. The outlet gas stream from the reactor was passed through dilute barium hydroxide solution for determination of C02. After the photocatalyzation process, the catalyst was filtered out, washed with distilled, deionized water, dried and weighed. Then, the catalyst recovery was calculated by applying the following relation: Recovery (%) = (W / W„) 100 -615- Where W = weight of the catalyst collected after the photocatalyzation process and W„ = weight of the catalyst before the process. To check for the activity loss of the aerogel catalyst, a series of experiments was carried out by mixing a dilute solution (522 ppm) of ferric cyanide with the aerogel catalyst (TiO,:Si02 molar ratio = 1:3.9), at the same conditions as before. At the end of the run, the reactor was rinsed and the aerogal was filtered out, washed and reused in subsequent runs. This procedure was repeated for three times at the same conditions. 3. RESULTS AND DISCUSSION 3.1. TiO,-SiO, Aerogel The alcogel structure contains more than 90% by volume fine pores containing alcohol.9,ao This alcohol must be removed to obtain aerogel. Because the radius of such pores is extremely small (nm to urn size) the surface tension at the interface between liquid and gas (vapor) is extremely high. To prevent damage to the gel structure due to these high interfacial forces, drying was done under supercritical conditions where the inter facial forces are minimum. After supercritical drying, dark and cracked monolithic aerogels were obtained. Then, heat treatment in air at about 550 C for several hours (about 5 hr) lead to converting the aerogel monoliths from dark and opaque to white and translucent. This is due to elimination of the organic residues." Figure 3 shows a photograph of two aerogel monoliths (after heat treatment) with different TiOa:Si02 molar ratios, prepared from equal volumes of alcogels. The effect of composition of the starting solutions, i.e., Ti02:Si02 molar ratio, on the shrinkage, apparent density and appearance of the produced aerogels are illustrated in Table 1. It is noticed that shrinkage and apparent density increased with decreasing Si02 content (see Fig. 3). This observed increase in shrinkage is due to the coalescence of gel particles which lead to decreasing the porosity as well as decreasing the -616- translucency.'''0 This can be explained by the fact that extremely porous and transparent silica aerogel, with surface area of about 1600 mVg, can be produced by supercritical drying, however, titania aerogel (in the form of anatase) can exhibit a surface area of about 120 mVg.' So, increasing SiO, content in the TiOj-SiO, aerogel could have a positive effect on both porosity and translucency of the produced aerogel. tfafiiii-|i>ritr. lnWr^ljkrf>rtifaiteiiiut*»iW..:ifti~A:h --*i-^i».;.(Aiiafc*j*r*&^ Fig. 3. Two aerogel monoliths after heat treatment [Aerogel 1 - with TiOjiSiO, molar ratio = 1:3.9] [Aerogel 3 - with TiO,:Si02 molar ratio = Is 1.3] Table 1. Effect of TiO,:Si02 ratio on shrinkage, density and appearance of TiO,-SiO, aerogel. Ti02:SiOj Shrinkage (%) Apparent Appearance molar ratio density Linear Volume (g/cm3) 1 : 3.9 25 59 0.27 Translucent 1 : 2.6 35 67 0.39 < the above 1 : 1.3 41 79 0.56 «the above -617- 3.2. Photocatalvtic Oxidation of CN~ Different forms of Tio2 semiconductor, used in this study as photocatalysts, are given in Table2. Results obtained from the photocataltic oxidation of CN" at these different forms of catalysts are illustrated in Fig. 4. Table 2. Photocatalysts used in the oxidation of CN". (T is Ti02, S is SiOj and G.S. is Grain Size) Photocatalyst Characteristics Aerogel 1 T:S molar ratio=l:3.9, G.S. < 3 mm Aerogel 2 T:S molar ratio=l:2.6, G.S. < 3 mm Aerogel 3 T:S molar ratio=l:1.3, G.S. < 3 nun Anatase T > 98%, Powder, G.S. < 1 urn Rutile T = 93.3%, Powder, G.S. < 200 urn C3 o> e u 2 4 Illumination Time, hr. Fig. 4. Photocatalytic oxidation of CN" in ferric cyanide solution (1000 ppm). [O Aerogel 1, but without illumination] These results have shown that in the presence of aerogel 1 (Ti02:Si02 molar ratio = 1:3.9), about 43% of ferric cyanide was removed after illumination time of 2 hr, and after 6 hr illumination time the removal reached about 94%. However, photocatalytic oxidation of CN" in the presence of anatase powder appears to be slightly lower than that in the presence of aerogel 1. This could be due to the higher surface area of aerogel matrices (several hundred mVg) in comparizon with the anatase powder. As shown in Fig. 4, the photocatalytic activity of the aerogel increases with increasing Si02 content, which could be attributed to increasing porosity as well as translucency (see Table 1). The higher efficiency of anatase powder than rutile form (Fig. 4) was in part explained by possible trapping of electrons in the anatase form at the surface where they can more easily react with oxygen.4 In addition, the higher particle size, i.e., lower surface area, of rutile powder in comparizon with anatase powder (Table 2) had a negative effect on the photocatalytic activity. Results of control tests with irradiated solution in the absence of Ti02 or Ti02(aerogel I) containing solution in the absence of irradiation showed that no or very little oxidation of CN" occured under these conditions (Fig. 4). In all experiments a white precipitate was observed when the outlet gas stream was passed through the dilute barium hydroxide solution, confirming that CN" was finally converted into C02 and N2. The estimated recoveries (%) of different forms of Ti02 catalyst, after the treatment process, showed that the aerogel catalysts were comletely recovered. However, it was difficult to recover all the Ti02 powder used, e.g., about 16 % of anatase powder and about 9% of rutile powder were contaminated with the effluent stream after filteration (see Fig. 5). Results obtained from the series of runs to check for activity loss of the aerogel catalyst are illustrated in Fig. 6. No significant change in the photocatalytic activity of the aerogel was observed. Therefore, Ti02-Si02 aerogel could be -619- used effectively several times for photocatalytic oxidation of CN" in aqueous solutions without activity loss. (a) (b) ,-,w.-~~.. - ,. n-^CTIWm^-. ,-,., M ^ |||M| 2 *-= L 2 -^ (c) S—=i ,' -•Mi-iT >att—MMM«ni» Ail • Fig. 5. Different forms of TiOj catalyst used in the photo-oxidation of CN". (a) Aerogel 1, granules (Grain size < 3mm) (b) Anatase powder (Particle diam. < 1 urn) (c) Rutile powder (Particle diam. < 200 um) -620- 100 90 4L Fresh Aerogel 80- Second Use ^ 1 3 Illumination Time, hr. Fig. 6. Reuse of aerogel 1 in the oxidation of CN" in ferric cyanide solution (522 ppm). CONCLUSIONS Increasing Si02 content in the Ti02-SiO2 aerogel results in aerogel matrices with lower shrinkage, lower apparent density (i.e., higher surface area) and higher translucency. This leads to higher photocatalytic activity for oxidation of cyanide in aqueous solutions. In general, TiO2-Si02 aerogel is expected to achieve superior performance to TiOa powder and colloidal due to: (a) Increased surface area and porosity leading to improved reaction kinetics and extent, (b) Increased translucency due to presence of silica (titania particles are opaque) leading to improved utilization of light energy, and (c) Minimized loss of catalyst, as aerogel monolith and granules are easier to handle and recover for reuse than powders or colloids. -621- ACKNOWLEDGMENTS The author is very grateful to "Department of Materials Science and Engineering, Ohio State University" and "TAASI Laboratory", Columbus, Ohio, U.S.A., for lab facilities to carry out this investigation. REFERENCES 1. S.Q. Hassan, M.P. Vitello, M.J. Kupfele and D.W. Grosse; J. Air Waste Manage. Assoc, 41, (1991), pp. 710-715. 2. D.W. Grosse; J. of the Air Pollution Control Association; 36, (1986), pp. 603-614. 3. M.S. Chandrasekharaiah and J.L. Margrave; Proceedings of Waste Stream Minimization and Utilization, Vol. 2 - Industrial Liquid and Gaseous Waste Processing, V.E. Lee, ed., Austin, Texas, April, (1993), pp. 5.1-5.6. 4. S.N. Frank and A.J. Bard; J. of Phys. Chem., 81, (1977), pp. 1484-1488. 5. S.N. Frank and A.J. Bard; J. of Am. Chem. Soc, 99, (19770, pp. 303-304. 6. D.F. Ollis; Environ. Sci. Technol., 19, (1985), pp. 480-484. 7. R. Phillips; Sources and Applications of Ultraviolet Radiation, Academic Press, New York, (1983), p. 47. 8. N. Serpone and E. Pelizzetti,eds.; Photo- catalysis, Fundamentals and Applications, John Wiley & Sons, Canada, (1989), p. 124. 9. S.J. Teichner; in Aerogels, J. Fricke, ed., Springer-Verlag, Germany, (1986), pp.22-30. 10.C.J. Brinker and G.W. Scherer; Sol-Gel Science, Academic Press, London, (1990). ll.B.E. Yoldas; J. of Non-Crystalline Solids, 38 & 39, (1980), pp. 81-86. 12.A.I. Kingon and J.B. Clark; j. Am. Ceram. Soc, 66, (1983), pp. 256-260. 13.D.M. Krol and J.G. Vanlierop; J. of Non- Crystalline Solids, 63, (1984), pp. 131-144. 14.T. Woignier, J. Phalppou and J. Zarzycki; J. of Non-Crystalline Solids, 63, (1984), pp. 117-130 -622- First International Spring School & Symposium on Advances in Materials Science (SAMS 94) 1520 March 1994 Measurement of The Thermal Properties of Co(N03)2 and Acetamlde by a Modified AC-Heated Wire Technique By S.R. Alalia, G.A. Attia, M.M. El-Naggar, T.A. El-Sayed A.A. El-Sharkawy ,N. K.Mina , M. Kenawy Department of Physics, Faculty of Education, Fayoum Branch, Cairo University, Fayoum, Egypt. Abstract: A modified AC-heated wire technique have been used for the measurement of thermal activity (b), thermal diffusivity (a), volumetric heat capacity (pCp) and thermal conductivity (X) of nitrate salt Co(N03)2 in the temperature range 30-120 °C and Acetamide CH3CONH2 in the temperature range 50-90 °C . DSC investigations have been also performed for both substances. The obtained results for Co(N03>2 show that (a) slightly decreases, (pCp) increases but X remains constant in the measured temperature range.For Acetamide, (a) slightly decreases, (pCp) increases, and (X.) increases with temperature, but all of them do not change appreciably on melting. The obtained values for Xfor both substances show that the main mechanism of heat transport is due to phonons as expected. The upper limit of the investigated temperature ranges was dictated by the occurence of unbalance of the measuring system and disappearance of the third harmonic signal. This observation requires further investigations to find out the reason behind it, specially because the same phenomenon has been observed for some other salts. * Dept. of Physics, Faculty of Science, Al-Azhar University, Cairo, Egypt. ** Dept. of Physics, Faculty of Science, Ain Shams University, Cairo, Egypt. •** Dept. of Physics, Faculty of Women, Ain Shams University, Cairo, Egypt. -623- Introduction Nitrate salts have many technological applications, specially in energy and solar engineering [1-5]. Measurement of their thermal properties: thermal activity (b), thermal diffusivity (a), volumetric heat capacity (pCp) and thermal conductivity coefficients (X) of these materials is essential for their proper use. From the scientific point of view, investigation of the mentioned properties helps the understanding of their stucture and the mechanism of heat transport in them. Experimental The AC-heated wire technique [6] was applied for the measurement of b,a, (pCp) and Xfor these substances. A platinum wire of The investigated Co(N03)2 and CH3CONH2 salts in powder form were melted in the experimental ceramic(and stainless steel) experimental cell in a high temperature furance. The platinum wire, fixed by means of two stainless steel electrodes in the cell lid, was then quickly immersed in the molten substance and then the cell was let to cool slowly. Then the cell, tightly sealed by means of silicon rubber sealant, was connected to the measuring system and several runs were carried out during the heating and cooling cycles. DSC investigations have been performed for both substances with an apparatus SHIMADZU 50 with resolution 0.01 °C. -624- Results and Discussion Co(N03)2 In figure 1 are given the obtained results for a 20.2 mg sample of Co(N03)2 in an Al crucible, at a heating rate 10 deg/min. As seen Co(N03)2 melts 59.6 °C with latent heat of 97.15 J/g for heating and 83.63 j/g for cooling. This fact indicates that Co(N03)2 solidifies in an another form, and this needs further investigation. In table I are given the obtained results of the measurement of thermal properties: thermal activity (b), thermal diffusivity (a), volumetric heat capacity (pCp) and thermal conductivity A. in the temperature range 35-120 °C. These results show that b decrease by 6 % in the measured temperature range, (a) and X they remain almost constant. Plotting these results show a pronounced peak at melting for b, a and (pCp) but not for X. CH3CONH2 In figure 2 are given the obtained DSC results for the investigation of a 19.7 mg sample (chemically pure Merck reagent) in Ai crucible. As seen Acetamide starts melting at 79 °C and has a peak at 89.9 °C with complete melting at t =120 °C,with latent heat of meltig 6122.4J/g and 3073.87 J/g for solidification. In table II are given the obtained experimental results for the mentioned thermal properties of acetamide. As seen b decreases monotonically till melting starts .while (a) remains constant. X and pep decrease up to 60 °C and again remain constant . Only two points have been obtained above the melting peak at t=82 °C. The upper limits of the investigated temperature ranges 116 °C for Co(N03)2 and 92 °C for CH3CONH2 were -625- THERMAL ANALYSIS DATA DATE 94/06/21 FILE NAME TAK106B.400 [DSC]- M T I ! >£>StfUKG CCHJITIOKS 10.0 SAJ4t£ NAME C0IN03I2 4.0 SAMPLE SIZE 20.200 S3 [nw/mln] SAHPLINS 1ST 3.0 tao ACQ. DATE 94/06/16 COMCKT 3rd Run (5C/«) tCATXNG PROGRAM RATE TEW TIlC t S.O 100.0 0 to 59 • 100 [c] Thareal Figure (1) DSC results for Co(N03)2 (heatirtg) THERMAL ANALYSIS DATA OATC *4/07/1t FILE NAME TAK108 .000 rosci n IC*SUR:S8 eeroltiOMS [•V] (Cl tw%e we *crr«;ce 52.0 aoo_i tMdS SIZE 19.709 Kg IMPIXNS XNT a.e ••• *C«. 0»Tf 14/07/31 COOCKT MUTXN8 PBCCa»« fUTC ItnP TX« I 19.0 128.0 t I -10.0 40.8 I -«s.o •J 0.00 -I Figure (2 a) OSC cesutls for AcctamK* (heating & cooling) Yt.oe uuil ItltTMl THERMAL ANALYSIS DATA DAtE 94/07/19 FILE NAME TAK107A.00Q "" [DSC] T—r -i—r —1—r -i—|—i—i—r T 1 I 1 T fCAsuniN'j COIOITJONS SA>»Pl.e NAHE ACCTAMICC SAMPLE SIZE J7.500 *} SAWLINfJ INT 3.0 a«a AO. DATE 9V07/19 ccworf reATIM<; PPCiflAM PATE TDf TI*E 1 19.0 130.0 0 89.9C I i t I I I I 1 L. L ' ' ' ' I ' ' ' i I 59 100 153 230 [C] SHIXAOZU Figure (2 b) DSC results for Acetsmide (heating) Tabic (1)Thmiiiil propertiesoff 'oONOJ, N»»'-' in1 C) ] axUFOn' *) ] Xiw/mC) [ ~,K(\\»"III'C) I 35.R9 1,151 1.43 0.42 2.90X.952 40.f>X I.07K 1.5 0.42 2.77X.XIK 55.07 1,176 1.23 0.41 3.350.954 5X.37 1,209 1.16 0.41 3,551,121 63.43 1,155 1.29 0.41 3,221,996 6X.45 1,096 1.45 0.42 2.X8U.251 ftV.X 1,0X7 I.4tXo 0.4«.•#.2; 2.X26.X09 104.4 1.0X0 1.5 j 0.42 2,7X9.421 115.4 1,046 1.46 j 0.42 2.X57.1X3 116.4 1,090 1.47 0.42 2.X45.443 .1 -629- TuMtf (II) Tlu-rmul pro|n-i lies of Acefamiilo b(v,s,J ni'C) a N U)'{m> s) Aiw/mC) pctwsim'C) 51.97 677 1.89 0.29 1,557,962 54.7 CU.J 1.17 0.21 1,820,(171 59.4') 494.09 2.38 0.24 1.016,949 60.7 5117.24 2.17 0.24 1,088,938 62.2 520.8 1.99 0.23 1.168.872 66.5 5UK.8 2.15 0.24 1,099,587 73.6 538.8 1.78 0.23 1,277,999 77.X 442..1S 2.14 0.2 1.080,586 79.5 470.8 2.7 X 0.25 892.X51.4 88.53 4V7.64 1.77 11.21 1.208,365 90.99 465.5 2.9 0.25 806,903.7 -630- dictated by the occurence of an unbalance of the measuring system and the disappearance of the third harmonic signal. This unbalance has been observed for other salts at different temperatures and requires further investigation to find out the reason behind it. Conclusion Further investigations of the thermal, electrical and other properties of nitrate salts and Acetamide at higher temoperatures to follow the structural changes in this class of chemical compounds. Nitrate salts are not simple as it was thought before. References: 1. Hirumichi O..Rev. Sci. Instrum., 61,10, (1990), pp 2645-2649. 2. Kato Y. et al, High temp.-High press., 15(1983), pp 191-198. 3. Araki N. et al, Int. J. Thermophys., 9,6, (1988), pp. 1071-1080. 4. Nicolic R. et al, Solar Energy Materials and Solar Cells, 28(1992)pp.59-69. 5. Asao Y. et al, J. of the Phys. Soc. of Japan, 17,3, (1962), pp. 442-446. 6. Atalla S. R. et al, Int. J. Thermophys., 2, 155, (1981). -631- First International Spring School & Symposium on Advances in Material* Science (SAMS 94) 15-20 March 1994 Measurement of The Thermal Properties of AgN03 and Cd(N03>2 by a Modified AC- Heated Wire Technique S.R. Atalla, M.M. El-Naggar, T.A. El-Sayed A.A. EI-Sharkawy*,N. K.Mina**, G.A. Gamal, andAXNagat*** Department of Physics, Faculty of Education, Fayoum Branch, Cairo University, Fayoum, Egypt. Abstract: A modified AC-heated wire technique was applied for the simultaneous measurement of thermal activity (b), thermal diffusivity (a), volumetric heat capacity (pCp) and thermal conductivity (X) coefficients of nitrate salts in the temperature ranges 50-190 for AgN03 and 35-75 °C for Cd(N03)2. DSC investigations have been also performed for both salts. The obtained results for AgN03 show a monotonic decrease of b,a and X, and an increase in (pCp). For Cd(N03)2, an increase in b,a, X and (pCp) was observed with pronounced peaks at melting. The upper limit of the investigated temperature ranges was dictated by the observed unbalance of the measuring system. This phenomenon has been observed for other salts and needs further investigations to find out the reason behind it. Dept. of Physics .Faculty of Science, Al-Azha University, Cairo, Egypt. ** Dept. of Physics, Faculty of Science, Am Shams University, Cairo, Egypt. *** Dept. of Physics, Faculty of Science, Kena Branch, Assiut University, Kena, Egypt. -633- Introduction Nitrate salts are an important class of substances that have many technological applications, specially as solar and thermal energy storing media [1],[2]. Measurement of their thermal properties; thermal activity (b), thermal diffusivity (a), volumetric heat capacity (pCp) and thermal conductivity {X), is essential for their proper use. This also helps the understanding of their stucture and mechanism of heat transport in them [3], [4], Experimental The AC-heated wire technique [5] was used for the measurement of the mentioned properties. A platinum wire of <|>50 ^m, immersed in the investigated medium, was heated by an AC- current of 0.25-0.30 A(50 Hz). Measurement of the applied power and the amplitude of the third harmonic of the temperature oscillations of the wire, enables the simultaneous measurement of b,a, (pCp) and X for the investigated medium. The investigated AgN03 and Cd(NC>3)2 salts in powder form were melted in a ceramic(and stainless steel) experimental cell in a high temperature furance. The platinum wire, fixed to the lid of the cell by means of two stainless steel electrodes, was then quickly immersed in the melt. After that, the cell was sealed by means of silicon rubber sealant. The experimental cell was then connected to the measuring system and several runs were carried out during the heating and cooling cycles. DSC investigations have been performed for both salts with an apparatus SHIMADZU 50 with resolution 0.01 °C. -634- Results and Discussion AgNO, In figure 1 are given the obtained results of the DSC investigation of AgN03 specimen (18 mg), in an AI crucible. The sample melts 216.9 °C. with latent heat of 62.28 J/g. As seen at T=169.7 °C an endothermic transition takes place. This transition was observed by several authors [6] during the investigation of dielectric properties of this salt. This is called ll-l transition, from orthorombic structure to trigonal. During cooling this transition did not take place, and the latent heat of solidification was 60.8 J/g. This particular behaviour needs further investigation to find out the reason behind it In table I are given the obtained results for b, a, X and pCp. As seen there is a monotonic increase in the values of b and (pCp), but the values of a and X decrease in the investigated temperature range. The behaviour of (pCp) show that the investigated temperature range is above the Debye temperature. The values of X show that the main mechanism of heat transport is due to phonons. The role of other mechanisms needs further investigation. Cd(N03)2 In figure 2 are shown the obtained results of the DSC investigation of Cd(N03)2 (17.9 mg) in AI crucible. These results show that the specimen starts melting at 55.6 °C and melts completely at 65.1 °C. The most striking feature of this salt is that it does not solidify on cooling Gil the room temperature, it remains liquid. This phenomenon needs further investigation specially because it was also observed for Pb(N03)2. In table II are given the obtained results for b, a, X and pCp for this salt All the results show an increase with a pronounced peak at 58°C. These results also show that the investigated temperature range is above the Debye temperature for this substance as expected. Concerning the thermal conductivity results, they show that the main mechanism of heat transport is mainly due to phonons. -635- THERMAL ANALYSIS DATA tUTT 94/04/24 FILE NAME TAK104 .000 [DSC] »CASUUNO cocmora r-wi auvLE HA>C 4aMn 20.0 SAX>1£ SIZE 18.MO ag 3AVR.IHQ IHT 3.0 ••• ACQ. DATE 94/04/tl COHCHT AL.OW.JD-W HZ fOTIMJ PROGRAM RATE TEM» T»€ I 10.0 300.0 0 3 -10.0 40.0 0 OS ON C»ln| Thiraal Figure (1 a) DSC results for AgN03 (heating & cooling) THERMAL ANALYSIS DATA OATC 94/04/24 L E E Q A O 7 , ,"^ , ,™<*, I ;°P . •.. • ^fl tCiSURIM CCMUTIONS (•Ml •A>ct.E Hue Aflicn 10.0 8AH>1£ SIZE 18.000 aa SU4X.IM IKT 3.0 ••• 4CO. DATE 84/04/OS COVCNT 4..BM.ID tCATIM PRCGAAM RATE TtM» TI« 1 10.0 280.0 0 ^1 iis.ac -30.0 I..I.I I I , I I I I, I 1 t I I I I I I I I I I t I I I till SO too tso 200 250 300 CCJ Tharul Figura (1 b) DSC results for AgN03 (boating) THERMAL ANALYSIS DATA 0»rC AVS4/S4 TAK103A.000 ICASUMMI CBClTloa •«*>Ut lizt 17.308 »Q •«*«n.tia IKT J.« ••• »0». BATE H/94/09 CMGff AL.BM.IO tCAtlNO fKXMM MTt tD* TIM I 9.9 IBS.* 0 • -•..• ao.t • oo -30.S _ O.09 M.Ot (ami Th*r*l Figura (2 a) DSC results for Cd^NO^ (heating & cooling) THERMAL ANALYSIS DATA DATE 04/04/24 rILE NAME TAK103 .000 >CASUUNO CDOITIOa IAHAJE w>e ot oral a •AMU SIZE 17.900 as 0A>4t.IM INT 3.0 1*9 AC8. DATE «4/04V!a CCMCMT «..Cn-CIO-(9-/.lnl tCATIM POOOUM BATE ro«> n>e i 9.: i».o o -30.0 _ toe (CI Th*ra«l Figure (2 b) DSC results for Cd(M03>2 (heating) Table l(l') l>(vvs"/iu'C) a x IO'OII'/S) A.(W/lllC) ,K(US/III'C) 50.4 658.7 2.01 0.3 1.453,474 57.23 | 716.7 1.61 0.29 1,792,228 64.5 ; 669.1 1.95 0.3 1,509,561 6X.X ' 6X2.19 I.XG 293 1,5X3,575 79.2 j 692.6 1.77 0.29 1,644,661 91.5 719.6 1.59 0.29 1,807,111 97 716.3 1.43 0.29 1,974,655 120.7 717.5 1.6 0.29 1,793,726 1.10.02 75-1.51 1.38 0.28 2,027.799 I.W.7 731.6 1.51 0.28 1,8X1,2X7 l-l'W 745.1 1.43 0.28 1,967,136 161.1 699.7 1.72 0.2«) 1,6X6,584 165.6 513.6 2.08 0.23 1.125,974 171.6 337.64 2.5 0.17 675,443.2 ISX 3')».«v 0.15 1,019,353 L..._ ~" Z1EZ -640- Tabic (H) Tlicmuil iiropcrtu's ol'('tl(NO,)2 J KC) l)(»v»'-"m'C) a x I07(III'/.H) A(W/IHC) pc(\vs'ni C) 37.63 XI »6.l)5 1.14 0.27 2.378,220 40.2 735.95 1.49 0.28 2,334,879 41.5 846.17 1.36 0.31 2,2K«V>52 4.1.23 K51.7 1.74 O.J5 2,030,025 50.6 X67.59 1.66 0.35 1128.995 5X 25.3.4 7.4 0.22 293,371.5 59.6 950.2 1.31 0.34 2,624,125 60.4X 313.9 3.27 0.1X 549,017.4 61.6 X92.XX 1.6X 0.35 2,102,169 63.6 X15.29 1.77 0.36 2,003,363 61.7 X29.14 1.87 0.36 1,914,340 67.29 9X4.17 1.52 0.38 2,519,110 6X.7 977.9 1.44 0.37 2,5X2,139 71.31 926.4 1.4 0.35 2.476.X71 72.4 X60.0X 1.6S 0.35 2,120,563 73,1 X 922.X 1.41 | (1.35 | 2,455,137 j -641- The upper limit of the investigated temperature range was dictated by the fact that we could't achieve any balance of the measuring AC bridge above 170 °C for AgN03 and 75 °C for Cd(N03)2. This phenomenon needs further investigation because it has been also observed for NaN03lCo(N03)2 and Zn(N03)2 at different temperatures. Conclusion Further investigations of the thermal and electrical properties of nitrate salts must be carried out at higher temperatures to correlate their structural changes with their physical properties. The behaviour of these salts is much more complicated than it was thought. References: 1. Hirumichi O..Rev. Sci. Instrum., 61,10, (1990), pp 2645-2649. 2. Kato Y. et al, High temp.-High press., 15(1983), pp 191-198. 3. Araki N. et al, Int. J. Thermophys., 2,155, (1981). 4. Asao Y. et al, J. of the Phys. Soc. of Japan. 17, 3, (1962), pp. 442-446. 5. Atalla S. R. et al, Int. J. Thermophys., 2,155, (1981). 6. El-Kabbany et al, Phys. Stat. Solidi(a), 95,495 (1986); 95, 127, (1986), and 98, 987 (1986). -642- EGfco/lShtt First International Spring School & Symposium on Advances In Materials Science (SAMS 94) 15-20 March 1994- Evaluation Of The Efficiency Of Hydrophobic Agents By Using Neutron Radiography G. Scherpks Atominstitute of the Austrian Universities, Vienna, Austria N. Ashoub Atomic Energy Authority, Nuclear Research Center, Cairo, Egypt ADSTRACT • Neutron radiography can be applied to the building industry in all processes where water and other hydrogenous substances are present In this work the effect of different hydrophobic (water protective) agents on porous building material is studied by using the neutron radiography facility of the Atominstitut in Vienna. The time relationship of penetration of these substances into porous building material is measured. The efficiency of these materials against water penetration is tested. The volume concentration profile of a liquid substance in a sample is also evaluated. 1, Introduction Neutron radiography is a visual non-destructive method of testing, having some features in common with X-ray testing. Since the first reactor neutron radiographs in 1956 /5/ it has been used for many applications, among others in - the field of building industry. These applications are connected with non destructive detection of water or hydrogrenous substances in the matrix of porous material. Neutron radiography allows the investigation of moisture penetration and moisture distribution in building materials. In Austria a lot of hydrophobic -643- agents are available and for that in this work the advantages of neutron radiography are used to study the time relationship of penetration and the • efficiency of hydrophobic agents from Austrian market in porous building material. These substances are of great interest because the conservation of buildings (especially of historical monuments) becomes more and more important. The reason is the rising environmental pollution (acid rain, another aggressive chemical substances). The manufacture of a hydrophobic substance is based on silicon. During application it penetrates into the porous material. After this process the solvent evaporates and a very thin (in the range of molecules) film remains inside the pores. The surface tension in the pores is changed and water will be rejected IV. In this work the Standard Testing Method Of Moisture Analysis (STMA) is used. The volume concentration profile of a substance in a porous sample is also provided. 2. Experimental conditions The neutron radiography facility at the TRIGA Markll reactor in Vienna is used. It was designed for industrial application beside educational purpose. The principal characteristics of the facility are the following IVJ2I: Thermal neutron flux 1.9 x 10* cm'V R With this method 73/ it is possible to study the penetration of liquid in a porous sample. The STMA consists of following steps: • Radiograph of the dry, untreated sample - Radiographs of the sample in selected time intervals during the treatment with a liquid substance (to get knowledge about the behaviour of the liquid border inside the sample) - Radiograph after the treatment of the sample (to see the finally treated part of the sample) The radiographs are evaluated with a microdensitometer. The blackness (optical density D) along the sample is determined. Fig. 1 shows a sandstone sample which is soaked in water. The change in blackness signals the liquid border. The information obtained from STMA is - maximum penetration depth of the liquid border in the sample - penetration speed of the liquid border black: water inside the sample bright: dry part of the sample &4ft Fig. 1 /!/ 4. Penetration of hydrophobic agents into sandstone Sandstone samples are soaked in different hydrophobic substances (5 mm layer) -645- and the Standard Testing Method Of Moisture Analysis (STMA) is done as described above. The results are compared to the penetration of water into the sample. Tested substances: Furtisol S4 LM (firm Franz v. Furtenbach AG) Sikagard 70 (firm SIKA Plastiment Ges.m.b.H.) Cryl-HP (firm STO AG) Redisan (firm Terranova) ft / cm » .-* MjO Furtisol S4 Fig. 2: penetration depth h as function of time III 5. Efficiency of the hydrophobic agents The sandstone samples which were treated with hydrophobic agents are used for testing the efficiency of these substances. The samples are soaked in water with ; the opposite side as before and with the Standard Testing Method Of Moisture Analysis (STMA) the penetration of water into that part which is treatet with the hydrophobic agent can be investigated. Fig. 3 shows the penetration of water in the hydrophobic parts of the sandstone samples. -646- h / cm Sikjgjra 10 *MMM • • Furtitol U CM Q . — O "~ —"*•* Cttt'M* « IM Fiff. 3: penetration depth h of water in the treated part of the sample lit 6. Determination of the volume concentration profile Additional to the penetration speed of a liquid substance into a porous sample some knowledge about the volume concentration profile of the substance in the sample can be helpful. The basic equation for a soaked sample is X = e",s,•*ct,1«•"1<, T...neutron transmittance through a sample ^...macroscopic crossection of the dry sample S|i,u4l|...macroscopic crossection of the liquid substance d...thickness of the sample in beam direction c.volume concentration of the liquid in the sample The evaluation of the distribution of the macroscopic crossection 2 along a sample is as follows. The sample is radiographed together with a standard. Here the standard consists of six Cd sheets of different thickness (0.05-0.3 mm). The blackness (optical density D) along the Cd standard which is measured with a microdensitometer deliveres a calibration straight line (together with the known -647- neutron transmittance through a Cd sheet of certain thickness). From the distribution of the optical density along the sample one gets (from the calibration curve) the distribution of the neutron transmittance along the sample and therefore the macroscopic crossection I (from the relation T = e"). The result received is dependent on the used standard and therefore needs not to be in agreement with other literature. Following steps are necessary to get the volume concentration profile: - Radiograph of the dry, untreated sample together with a calibration standard to get the distribution of 2^, along the sample - Radiograph of a defined volume of the liquid substance together with a calibration standard to get Z,^,, - Radiograph of the soared sample together with a calibration standard to get the distibution of !„, s Ej, + cl^^^a along the sample From this equation one can calculate the volume concentration c of the liquid along the porous sample. This evaluation process is called Neutron Transmission Analysis (NTA) /4/. During one series the experimental conditions must be constant Here the volume concentration profile of a hydrophobic agent (Furtisol S4 LM) in a sandstone sample is studied. The sample is soaked in the 5 mm liquid layer for 2 min. The densitometric measurements were made at the Paul Scherrer Institut (PSI), Villigen, Switzerland. Fig. 4 shows the volume concentration c of Furtisol S4 LM along the treated sandstone sample. This result agrees with a weight measurement which was made on trial. Fig. 4 IV • 10 ItaflkafMaifk/cM 7. Results and conclusion The time relationship of penetration of hydrophobic agents and water into sandstone samples were measured. The penetration speed of water into the sample is higher than that of the other substances (Furtisol S4 LM, Sikagard 70, Cryl-HP, Redisan). A maximum penetration depth cannot be determined, see Fig. 2. All tested substances show efTiciency. It can be seen that the capillarity elevation of water into the treated part of the sample is very slow compared with that into the untreated one. At a certain level the water border is stopped. Fig. 5 shows a little difference in the efTiciency of the tested substances. Neutron radiography is also a very convenient method for studying the concentration profile of a liquid in a sample. It can be assumed that the hydrophobic agents available on Austrian market are of good quality. 8. References IV G. Scherpke: Untersuchung von Feuchtigkeitstransport in BaustofTen mittels Neutronenradiographie, Diploma work, TU Wien, Austria, 1993 /2/T. Buchberger: Neutronenradiographische Untersuchungen an Tantal-Tritium- Helium-Systemen, Dissertation, TU Wien, Austria, 1986 121 F. Peterka, H. Bock, A. H. Pleinert: Neutron Radiography Standard Testing Method for the Moisture Analysis in Building Materials l\I Z. Hridlicka, F. Petcrka: Principles of Neutron Transmission Analysis 151 J. Thewlis, R. T. P. Derbyshire, A E. R. E. M/TN 37 (1956) -649- EGfSMMS First International Spring School A. Symposium in Advances in Material Science (SAMS 94) Cairo, 15-20 March 1994 NEUTRON DIFFRACTION AND MOSSBAUER MEASUREMENTS ON GALLIUM SUBSTITUTED ZINEFERRITE M.K. Fayek, F.M. Saycd Ahmed, H. Abou-Hela!, M.F. Mostafa and M. Kaiser. Dept. of Reactor & Neutron Physics, Nuclear Research Centre, Atomic Energy Authority, Cairo, Egypt. Abstract Three powder samples of Zn Gax Fe2_x04 (0 £ x £ 0.6) were syntliesised by using a ceramic sintering technique. X-ray powder diffraction pattern were obtained for these compounds and confirmed the presence of single-phase spinel structure with no evidence of impurities. Neutron diffraction measurements were pcrfonncd for the polycrysialiine powder Zn Gag -joi Fej AggO^ sample. The atomic coordinates and cation distribution were determined based on X-ray and neutron diffraction patterns. The cation distribution formula is then: ,*. 2+ „ 3+ „ „ 2+ „ 3+ „ 3+ , ,. (Zn Ga ) [Zn Ga Fe ] 0 0.74 0.26 A A 0.26 0.07 1.67 u n 4 4 MOssbauer effect patterns were obtained and revealed a paramagnetic behaviour at room temperature. The patterns indicated that no iron is left at the tctrahedral site confirming the obtained neutron data. The spectra were successfully analyzed and the different MCssbaucr effee parameters were deduced. 1. Introduction Ferrites are semiconductors containing one of the trivalent ions as major metal component. Many of its properties strongly depend on their exact chemical composition and microscopic physical structure. Cubic mixed fcrrites arc the most important ferroxcubes. -651- Its characteristic is that the physical properties changes completely or partially with changing temperature. Ferroxcubcs, have many uses in high frequency electrical and electronic devices such as transformers, dynamos, amplifiers and as a memory clement in computers. Zinc fcrrite together with a large number of other spinels were investigated (1-3). Recently, there has been a considerable revival of interest in the properties of magnetically disordered ferrhes. In particular, there have been several neutron diffraction and Mflssbauer effects studies of fcrrimagnclcs and aiitifcrronngncts of the cubic spinel fcrrites (4-8). In this paper, we report the results of X-ray diffraction, neutron diffraction and Mbssbaucr effect studies on powder samples of Zn Gax Fe2.x O4. 2. Sample Preparation: The investigated mixed zinc, gallium fcrrites were synthesised using die usual ceramic sintering process (9). Powder forms of stohiomctric ZnO,Ga2C>3 and aFe2C>3 were mixed in the appropriate equivalent weights and annealed in air at 1000°C for (48 h). The compound were pressed into pellets and annealed again at 1150°C (48h) then followed by slow cooling. X-ray powder diffraction patterns were obtained for die compounds, using Coka radiation, fig. 1 shows single phase of die spinel type with no extra lines due to impurities. The lattice parameter were found to be = 8.4594,8.3826 and 8.3664 for X=0,0.333 and 0.666 respectively. 3. Experimental: The neutron diffraction pattern of the prepared powder sample ZnGan .mFei (MAOA has been obtained at room temperature fig.2, using a -652- XJOO o o «* *n ~O *n i b£. A L 1 -A. X:033 I_L X*0 66 Vj_ J-^-i 10 li 80 70 26 ; I'ifl- I: X-my tlilTniciitui penicm of Xii (!:ix I'O-x Q4 l crriie .sysicm. -653- s 16500- 3 X= 1.08A H500- 12500- £10500' at 8500 6500 4500- 2500 —i 1 1 1 1 1 1 1 1 1 10 15 20 25 30 35 40 45 50 55 60 20 Fig. 2: Neulron dirrraction pattern of the compound Zn Gao.333 Fci.gg^ O4 at room temperature. -654- neutron diffractomctcr placed at one of the beam holes of the ETRRj at a power of 2MW. The monochromatic neutron wave length was 1.0825 A". AH measurements were repeated twice. The obtained measurements did not give any appreciable difference. The data shown in Tig. 2 is just one of them. The observed intensities of retlections were obtained in the usual way. Mbssbaucr resonant absorption spectra (ME) of the studied compounds were taken at room temperature fig. 5. The source used was 50 mCi Co diffused in pd matrix. The solid lines through the data points are the results of the least squares fit to the data. The determined ME parameters arc collected in table (3). 4. Results and Discussion: 1. Crystallograghic Parameters: The crystal structure of the prepared compounds Zn Gax Fc2.x O4 are isomorphous with the mineral spinel. The symmetry is cubic, space group Pd3m with eight formula units in the cell. The metal ions arc in special positions without parameter and arc distributed in 8 (a) tclrahedral A-sitc and 16(d) octahedral D-siie. The oxygen are in the general positions 32(e) and form a cubic close-packing. The crystallographic parameters to be determined were: i) The coordinates for oxygen atoms in the unit cell (u.u.u) ii) The degree of inversion (i) The two paramctcs were determined using neutron patterns and comparing their observed and calculated intensities. Coherent neutron scattering amplitudes as reported in Bacon (10) were used. The reliability factor R was minimized by a systematic -655- Table (3): Mossbaucr Effect Parameter for Zn Gax Fe2.x04 at room temperature Isomer shift Quadriipole Compound (in m s" ) Splitting (m m s ) Zn FC2 O4 0.523 0.195 ZnG.To.33Fc16604 0.513 0.205 Z'>Gao.66FcL33 04 0.431 0.267 IZrror ±0.02 ±0.02 -656- 0-115- 0.095- u £a 0.075- •g \= 0.055 Zn - B-sites No. i 0.035- 1 0.00 2 043 3 o.oe 4 0.09 s 0.13 0.015-1 1 1 r 0.365 0.375 0.3B5 0.395 0.A05 Oxygen Parameter U Fig. 3: Minimization of the reliability factor R with the characteristic parameter U and i for Zn Gag 333 Fej 6^5 O 4 -657- .^ t I, "3 * 3 "• S. fl 4? 601 ^J7 -658- .10" 8.0866 7.9308 "" • r" 7.7 678 7.6048 x=o.o 7.4 418 J- 6-7 0 6 «105 6.668 w» 6.632 c: 2 6.592 J- >. * 0.333 6-5 54 4.089 4-0704 , , .^'J'' ^^•^^' }\ AT- ^ 4.0512 4.0331 X =0.666 4-0146 L. J__i I i L i • i _J -6-4-2 0 2 4 6 8 Velocity mm /sec /•'»*. 5.- Ml-: S|KIII;I l».r Hie Icniic sysicm X.nCJax l-Vi-x ()4 ai room icinpcniiurc. -659- variation of the above parameters. 2/JF0"JFc/ R= =— Where : J is the multiplicity factor, F0 is the observed structure factor, Pc is the calculated structure factor. The minimization of the reliability factor R against the oxygen parameter u is represented in fig.3. (R = 0.019). The Debye-Waller temperature factor was obtained 2 2 2 in the usual way from a plot of jFc/I0 against sin 0/X fig.4 A comparison between the observed and calculated structure factors after introducing Debye-Waller temperature factor is shown iii tabie (1), and the best agreement with the reliability facJ'V (R= 0.011) was obtained. This value of the reliability factor (R=0.011) is literally considered to be a measure for a good fitting. The proceeding analysis of the experimental results supports the structure of the ferrite Zn Fej ggg G°0.333O4 to be inverted spinel. Moreover the good agreement between the observed and calculated structure factors means a high precission in value of the oxygen parameter and reflects the proposed atomic distribution in the unit cell. Satisfactory agreement between observed and calculated squares of structure factors are obtained for an oxygen parameter u=0.383 and reliability factor R= 0.011. The comparison between observed and calculated intensities for Zn Gao.333 Fej^(P^ sample is given in table (1). The absence of any extra lines and reasonable low value of the reliability factor proves that the studied compound has no magnetic contamination and followjf -660- Table (1): Comparison between observed and calculated nuclear intensities for Zn 039333 ^1,666^4 at room temperature. hid iFj rf 111 148 191 220 555 456 311 4286 4286 222 999 954 400 5268 5268 331 42 61 422 785 820 333 \ 5193 5193 511/ 440 1786 1790 531 1484 1333 444 693 . 559 -661- the paramagnetic behaviour at room temperature. The X-ray and neutron diffraction results show that the galluim and zinc are distributed between (A) and [B] sites. The deduced formula could be written in the expressive form for neutron as: (Zn0 j4 Ga0 26)A [Znft 26 Ga0 07Fc, 67] B04 and for X-ray as: Zn Ga + Fc Ga Fc ( 0.730 0 187 6.083>A rZn0.271 0.146 *l.583l B°4 The given cation distributions in die (A) and [D] positions for X-ray and neutron experiments differ. This difference occurs because for X-ray diffraction the scattering power of an atom increases with die atomic number and decreases with the scattering angle while there is no variation between scattering power and atomic number of die scaltercr for neutrons. The interatomic distances and intcrbond angles in the unit cells of the present Zinc Gallium Fcrritcs are calculated using die obtained values of the cell and atomic parameters. The results are represented in table (2). 2- Mossbauer Effect Results : The Mflssbaucr effect spectra for the compounds under study at room temperature fig.5 show quadrupolc doublet with assymciry in the peak intensities. In order to get analysis for the ictrahcdrai and octahedral contributions, the full widdi at half maximum of die obtained quadrupolc doublets for all compounds ((KxS 0.666) was compared with that of a-Fc2C>3. It was found that its values arc less with those of a FC2O3. 2+ No line due to Fe in die tctrahcdral was observed due to large concentration of Zn in die tctrahcdral site, which menns dial Fc ions arc located specifically at the -662- Table (2): Interatomic distance and intcrbond angl For Zn Ga0>333 Fe1#6 For Tetrahedral For Octahedral 0 A-Olct =5.7 (12 times in A°) B - Otct = 3.7 (9 times in A ) °tct-°ict = 57 <2,imesinAC) °lct * °tct = 5-2 & limcs in A°> Otet-Otet-O.ct=71° Oict'Old-Otct =55° 0,et-OtcrA «=180<» 0,et-OiCt-B =82" 0 0,et-A -0,ct =0° Owt-B-Otet'SP -663- D-lattice sites. This conclusion is confirmed by the cation distribution formula obtained from the neutron diffraction results for the compound ZnGan.333 Fci gg^. It is clear from table (3) that the isomer shift values which is rclntive to iron metal for ferric ions lie between 0.431-0.523 which is the range characteristic of the trivalcnt high spin state of the iron (Fe + S= 5/2). Low temperature measurement arc naw in progress in order to determine ihc magnetic properties of the present samples. -664- References Craus M.L., Rczlcscu E., and Rczlcscu N.; phys. Sat. Sol. (a) 133,439 (1992). Palil H.D., Upadhyay R.V., Shamkuwar N.R. and Kulkami R.G.; Solid State Commun., 81, 1011(1992). Pannaparayil T., Komarncni S., Marantic R. and Dezsi I.; Hypcrfine Interactions, 54, 645 (1990). Fa-Shcn Li, DE-Shcng Y., llua-xin Z., Xin-fcn L. and Zheng Y.; Hyperfine Interactions, 41,525 (1988). Melzcr K.; Crystal Res. & Tcchnol., 17 (1982) /K3- K5 Bhargava S.C.; Mulder F.M„ Tliicl R.C. and Kulshrcshlha S.K.; Hypcrfine Interactions, 54,459 (1990). Fayek M.K., Bahgat A.A., Abbas Y.M. and Mobcrgt, J. Phys. C: Solid State phys., 15,2509 (1982). Fayek M.K., Arab Journal of Nuclear Science and Application 233-250, 15-1 (1982). Pointon A.J. and Robertson J.M.. Phil Mag. 12. 725 (1965). Bacon, C, Neutron Diffraction, Monographs on the physics and chemistry of Materials pergamon press (1962). -665- First International Spring School & Symposium in Aduances in Material Science (SAMS 94) Cairo, 15-20 March 1994 THE DETERMINATION OF CU/NI RATIO IN COPPER-NICKEL ALLOY BY NEUTRON CAPTURE GAMMA-RAY METHOD Waleed H. Abulfaraj, Samir Abdul-Majid & Osama Turkustani Nuclear Engineering Department, Faculty of Engineering King Abdulaziz University, P.O. Box 9027, Jeddah-21413 Saudi Arabia Abstract Copper-Nickel alloys have wide application in industry, such as in heat exchangers pipes of desalination plants. A method is needed for analyzing nondestructively the Cu/Ni ratio of a large number of samples regardless of their physical shape or surface conditions. In the neutron capture gamma-ray technique, applied in this work, neutrons emitted from a 241 Am-Be source are allowed to interact with the materials under consideration after being first slowed down by a paraffin wax moderator to increase the neutron capture probability. Following neutron capture, several characteristic gamma-rays are emitted from Cu and Ni elements and measured by a HPGe detector. The intensity of gamma-rays-of each element was found to be proportional to the number of atoms of that element on the alloy. The wieght ratio of each element could thus be determined accordingly. The system was found to detect a 1% change in Ni to Cu ratio, with a 0.1% detectable change at a higher neutron flux value. 1. Introduction The Cu-Ni alloys are used widely in industry. One important application is in heat exchangers for desalination -667- plants. The weight ratio of these two constituents of the alloy will determine the desired physical properties as well as the alloy's material commercial value. It is, therefore, important to determine any deviation from, the specified ratio, preferably in a non-detructive way. Testing and scanning all the materials under consideration, rather than just a few samples would have clear advantages. The sample analysis method currently used is not practical for testing and scanning all the materials under consideration. Moreover, it is slow and destructive. It is, therefore, desirable to develop a fast and relaible technique that tests, nondestructively, bulky samples. Neutron capture gamma-ray method is suggested as a practical testing method for such an application. In this mehtod the detection and measurements are contactless. Neither the signal emitter nor its receiver are in physical contact with the material tested. Therefore, measurement can be made on hot, unclean or unprepared surfaces. The results of the measurement are produced by electronic equipment with outputs in digital or analogue forms appropriate for display, recording or connecting to a control system. Moreover, on-line and in-situ measurement can be made without interupting the work system. This is particularly important if the alloy is part of a component used in a production plant; no plant shut down will then be necessary. In the neutron capture gamma rays, neutrons emitted from the source are allowed to be incident on the material under consideration after being slowed down by a moderator, some of them will be captured due to the (n, 7) reaction, followed immediately by the emission of gamma-rays which are characteristics of the absorbing elements. Fast neutrons may undergo inelastic collision due to the (n, n' y) reaction associated with the emission of the characteristic inelastic gamma-rays. The emitted gamma-rays can be detected by a HPGe detector. The intensity of the characteristics -668- gamma-rays emitted from each element is proportional to its total amount. The analysis of the spectrum will, therefore, give information on the alloy constituents. Recently Abdul-Majid et al^1'4' , have developed the neutron capture gamma-ray technique for the measurement of general corrosion, scale deposition, as well as for identifying the type of scale likely to be found at desalination or Chemical Plants. The main applications of the neutron capture gamma ray method remain in the field of mineral explorations15"71. Evans18', and Hassan191 used the technique for in-situ multielement analysis, while Chang and Yuan have used it for in-vivo elemental analysis'^1. 2. Experimental Work The arrangement for the neutron capture gamma ray method is shown in Fig.l. It consists of 1.11x105 MBq(3 Ci) 241Am-Be (half life is 433y) neutron source that emits 6.6x10 n/s with a tolerance of ±10%, surrounded by 8 cm paraffin to moderate fast neutrons in order to maximize the slow neutrons at the sample position and to increase the capture process. The alloys tested where type C 70600 of 90% Cu and 10% Ni and type C 71500 of 70% Cu and 30% Ni. The sample dimensions where 15x10x14 cm.. The gamma rays are detected by a HPGe detector (type Gem-10175 Eg&G ORTEC, USA) of 10% relative efficiency. The detector was separated from the moderator by 10 cm lead to stop y-rays coming directly from the source or from the moderator following capture processes. The MCA used was of 8192 channels (type 5608 EG&G ORTEC, USA) used wiht the other conventional electronic components. Energy calibration was performed by using several well known gamma rays. 3. Results Measurements were carried out first without any sample to find out background peaks, particularly that of Cu -669- because of the several copper wires around. The background counts for 6 hours collection time clearly revealed the Cu 7.636 MeV single escapre at 7.403 MeV and double escape at 6.892 MeV appeared, but no Ni peak was found. When the damDle was irradiated for the same period of time, the Cu 7.636 MeV and the 7.914 MeV full, single escape and double escape peaks were clear much higher than the background. The spectrum of Cu-Ni alloy type C 71500 is show in Fig. 2. The Ni 8999 full, single escape and double escape peaks were also very clear. The net counts for this peak vesus Ni ratio are shown in Fig. 3. The counts increased almost linearly with increase in Ni ratio, particularly the double escape peaks, which showed higher counts over the single escape and the full peak. The neutron flux was measured by indium soil activation and was found to be 4xl03 n/cm2.s, 10 where the neutron fluence was about 8.4x10 n/cnr\ This will give a sensitivity of about 7x10 counts per unit fluence change per 1% Ni ratio change. 4. Discussion and Conclusion The data of neutron capture gamma rays showed almost a linear relation between the intensity of the gamma-rays and Ni concentration. The double escape peak showed higher counts attributed to the detector efficiency. The neutron absorption cross sections of Cu and Ni have similar values. Therefore, the change in concentration will not change, significantly, the neutron flux within the sample. The sensitivity to elemental concentration in an alloy depends strongly on the absorption cross section of the element and, therefore, will not be the same for all elements. Nickel has a low thermal neutron absorption cross section of 4.43 barn. It is, therefore, not considered as a sensitive one. Considering that a difference in counts of two standard deviation is measurable with good confidence level, a change in Ni of 1% in the alloy is measurable. It is quite possible to use a neutron source that would give more than a two order -670- •tit. '-r..i;J Tig. 1 Set-up for neutron capture gamma ray teclmique. Fig. 2 Spectrwn of neutron capture fpMna rays emitted from Cu-Ni alloy type C71500. -671- 2000 n c 3 o 1000- %> -o Ni(I) -* Ni(s) t> Ni(d) Ml ratio V. Fig. 3 Net counts for 6 h of 0.999 MsV full, single escape ond double escape neutron capture gamma 7.s>s vs. \'i ratio in Cu-Ni alloy. -672- of magnitude higher thermal neutron flux. At such a higher flux a concentration change of 0.1% will, be measurable. A source like 252Cf that has softer neutron spectrum, and about 1000 times as many neutrons as that of 241Am-Be of the same activity, will give higher detection sensitivities for changes in alloy concentration. Other possible improvement can be achieved by using a higher efficiency HPGe detector, but the system would be more expensive. Such detectors could improve sensitivity by one order of magnitude. Improvement in the geometry is possible also by reducing the background gamma-rays coming directly from the source, and by surrounding the sample by a moderator in order to increase the slow neutron flux. One of the main advantages of this technique is that a sample of any shape can be tested, and not sample preparation is required. The technique is contactless and the signals are given in a convenient digital or analogue form. The neutron capture gamma-ray method is not applicable for the detection of any element. The detection will depend on the element cross section and its gamma emission characteristics. Moreover, the produced radioisotopes should not have long half, otherwise the sample would stay radioactive. The HPGe detector undergoes deterioration under fast neutron irradiation and will need annealing after a neutron fluence of about 10 n/cm . In this respect the n-type detector is preferred over p-type. References 1. Abdul-tnajid S. "Corrosion and Scale Measurement in Iron Pipes by Prompt Gamma-Ray Analysis" Nucl. Instr. Meth. 73, 398,1993. 2. Malki B., Melaibari A., and Abdul-Majid S. "Measurements of Organic Scale by Neutron Capture Gamma Ray and Neutron Moderation" Trans. Am. Nucl. Soc, 69, ISO, 1993. -673- 3. Abdul-Majid S. and Dawood U. "Iron Pipe Wall Thickness Measured by Neutron Capture Gamma Rays" Trans. Am. Nucl. Soc, 63, 136,1991. 4. Abdul-majid S. and Dawood U. "Neutron Capture Gamma Ray Technique for Scale Identification Inside 'Pipes" Desalination 75, 199,1989. 5. Glascock M.D., Coveney Jr. R.M., Tittle C.W., Gartner M.L. and Murphy R.D., "Geochemical Applications for Prompt Gamma Neutron Activation", Nucl. Instr. Meht. in Physics Research BIO 111,1042, 1985. 6. Wormald M.R. and Clayton C.G., "In-situ Analysis of Coal by Measurement of Neutron-Induced Prompt Gamma-rays", Int. J. App. Radiat. Isot. 34, 71,1983. 7. Underwood M.C. and Dyos C.J., "Inelastic Neutron Scatterin Reactions in Fluid Saturated Rock as Exploited in Oil Well Logging", Appl. Radiat. Isot. 37, 475, 1986. 8. Evans L.G., "In-situ Elemental Analysis Using Neutron-Capture Gamma-Ray Spectroscopy", Nucl. Instr. Meth. 193, 1982. 9. Hassan A.M., El-Kady A. and El-Exaley B., "A Prompt Gamma-Ray System for Elemental Analysis of Complex Samples", Nucl. Instr. Meth. 192, 595, 1982. 10. Chung C. and Yuanm L., "Determination of Ca, CI, N, and P by In Vivo Activation Using a Mobile Reactor Neutron Beam", Appl. Radiat. Isot. 39/9, 977, 1988. -'614- First International Spring School & Symposium on Aduances in Material- Science (SAMS 94) 15-20 March 1994 Mossbauer Investigation Of A Canted Spin Fe(IH System M.F.MOSTAFA , V..ryi.;,BDELKADER. A.S.ATALLAH .S.SOLIMAN DEPARTYMENT OF PHYSICS.FACULTY OF SGIENCE.UNIVERSITY OF CAIRO M.EL-ESAWI DEAF-rtTMENT OF CHEMISTRY.FACULTY OF SCIENCC.UNIVERSITY OF CAIRO AND R.EMRICK DEPARTMENT OF PHYSICS, UNIVERSITY OF ARIZONA.TUCSON.AZ 85721 ABSTRACT Mossbauer masurements of hexanediammoniumtetrachiorolenaie over the temperature range 78-320K lor powdered and single crystal absorbers show a transilion la the aniilerromagnetically ordered slate at 105 0 » 0.2K . The center shift shows the con.pound to have an electronic ground stale of 3d* 4s". The quadrupole splitting indicates that Fe(ll) exists in a distorted octahedral field and that the ground state is a doublet, with the principle axis oMhe' electric field gradient being negative. The recoiless traction, the center shift and the quadruple splitting reveal the existnece ol a slructrual phase change at 235 ± 5K.lhe low temperature phase having a Debye temperature ol 166K and (he high temperature phase has a Debye temperature of 15SK The compound undergoes long range order at -99K . the internal magnetic field makes an angle of 80*1* with the principle axis of the electric field gradient, and H„ has a value ol 187 « kOe at 78K. " INTRODUCTION HexanediammoniumFe(ll)CI4. belongs lo the series 'CHj)„ (UH}). MX,, .where n=1.2. . and M- Mn, Co, Fe.Cd.Cu and X= CI and/or Br. It crystallizes in a layered type structure whereby the melal ion is surrounded by it distorted octahedron of CI ions. The*.' octc^Rdra share corners and extend in an infinite two dimensional network (1-14). The octahedra are tiiieJ back and forth forming a "puckered structure " . These puckered layers are sandwitched between the organic ammonium layers and are linked to it via N-H...CI bonds. At room temperature . members of the above series with even n are monoclinic while those with odd n are orthorhombicflO) The crystal structure of some members of this series containing Mn and Cd have been determmed(3.10). however, for fe(ll). which is isomorphous to the Mn and Cd, only the structure of the n«3 compound has been sclved( 6). Several successive structural phase transitions have been reported In the Mn and Cd compounds ot this series between 77K ;nd 370 KfU).These transitions are connected to the dynamics of the alkyienediammonium chains mostly due to order-disorder transitions of the chains or to their conlormotional changes. However , studies by Kharizi et <•! (16) showed that not only the organic molecules which are responsible for the transitions .but that (he melal ions play an important role in this matter as well. Magetric susceptibiltiy measurements ol the s« • •» where Ms Fe(li) was carried out between liquid nitrogen temperature and room temperature, yet no evidence of structural phase changes was found (11*13) However, all compounds will.- n «i .4.5.6.and 7 showed short range magnetic ordering taking plac. •• T- 110K with long range ordering at T -100K which is associated with a ferromagnetic interact*- • •.; a result of the canting of the spin syslem.As compounds contain'*) Fe ions are excellent candidates for 'tossbauer effect technique, it is very senstive to any changes in the crystal field «inr/ bonding of tr.i r.>(ll) ion. thus the compound HexandiammoniumFe(ll)CI„heiiceforth HDFC. was prepared and its ME spectra collected between liquid nitrogen and room temperature . in an attempt to study structrual phase changes .if any, and also to acquire a better insight of its magnetic behaviour at lower temperatures. EXPERIMENTAL Crystals of HOFC compound were prepared u/ mixing equimolar ratios of 1.6-diaminohexanehydrochlorkle with lerrous chloride in acidified aqueous solution under Nitrogen atmosphere. Slow evaporation at 50°C resulted in plate-like crystals. The crystals were then washed wiln a mixiure of ether and ethanol and were dried in a vacuur desscator. -675- The Mossbauer absorbers v.;re prepared by using both single crystals and/or powdered crystals in Apiezone grease, place.' n a mylar sample hotter of 2.5 cm diameter pcovtoed wiirt an O-ring to prevent oxidation. A 20 mCi Co" in Rh source (at room temperature) was used. A standard transmission geometry setup with a constant acceleration electromechanical drive system is connected to an IBM computer lor data collection The velocity calibration was performed using both Ve toil and Na,(FeCN).6H,0 powedered crystals. A Jenkins cryostst was used for the low temperature studies . A temperature control unit with accuracy of <0.SK was used. Results and discussions: 1. Center shift and Recoiless fraction A representative Mossbauer spectrum of the powder sample between 112K shown in Rg(1a). The well spHt quadrupole doublet of o.E0= 2.19 mm/s at room temperature is typfcal of Fe(U) ion in a distorted octahedral field. The value of the center shift at 297K of 1.15i0.01 mm/s (referred to a Co in stainless steel source) corresponds to a 3d* 4s" electronic configuration using Oanon's calibrat«n{l6}. The Mossbauer center shift consists of two parts , a temperature independent part, isomer shift rf„ and a temperature, dependent part which is the second order Doppfer shift rfMW The latter results from the relativists motion of the emitting nuclei and is given , to a first approximation by: Bt^JdT * -<3/2)|li /M..C') (1) where M„ is the effective mass of the Fe atom and C is the velocity of light(17). The temperature variation of the center shift between 112K from which one obtains 6lT (region (a)) * 166 » 3 K and Grt ( 2. Quadrupole splitting: As mentioned above a well split Quadrupole doublet .occurs throughout the temperature range 112rVi"< 320K. The temperature dependence of the EFG (AE0) is shown in Fig(5). A gradual increase in AE<, is associated with the decrease of temperature down to -235K where a sudden rise in AE0 occurs, followed by another monotonous increase in AE„ dowri to T-112K. A sharp increase is found below that temperature. Anomalies in the variation of AEg with temperature could very likely be due to structrual phase transitions and /or to magnetic ordering. Th 676- ]toooe k )»»ooo h S Jjoooe 3O000 H ]4oo»o y I.•!»••».« l.l***Ot Jl. •>»••»• I.IS*<0» t.M»««o» K l.M««Ot i.m* Fig 1. Mossbauer spec'.a ?t room temperature for a powdered sample and a single crystal with gamma-ray perpendicular to the plane containing the FeCI6 octahedra. -677- CL92 BO 120 160 200 2£0 280 320 Temperotorr I K I •••g.2. Center Shift as function cf ter.*perature between 412K<*<320 K -678- -0.20 0.00 - ,-0.20 - -0-tO —-0.60 < -0.80 h -1.00 - 1.2 0 80 120 160 200 240 280 320 Temperature t K J Fig.3. Temperature variation of the normalized recoiless fraction in the temperature range 112K -679- onset of magetic ordering as the obtained Mossbauer spectra start to undergo magnetic hyperfme splittings at 11 OK. This will be discussed below. As is well Known (18) the 'D state of Fe"2 splits in the presence of an octahedral crystalline field into a low lying triplet T„ and an excited, doublet E,. A Mragonsl ditlortion further splits the low lying doublet into either a ground state singlet ld„> and an excited stale doublet ld„>. ld„> or vice versa. The singlet state is associated with a positive sign for the mim component of the EFG (Vu) and the doublet is associated with the negative sign for Vu. In order to determine the sign of Vu and hence the direction o' the quadrupole splitting, one has to either apply an external magnetic field lo the powdered sampie or to use a singlo crystal with known orientation with »aspect to the gamma-rays In this experiment we have used a single crystal placed w.-!lh the gamma rays perpendiculji to the layrrs of FeC!t . The room temperature spectrum is shown in Fig.(1b). The relative intensities of the two peaks were found to '03 l,(low energy ) : l,(high energy) = 1.7 : 1 This agrees with 6 - 20 . * * 0.0 and n - 0.6 and confirms that V„ is negative, Iherfore suggesting a doublet ground state idu>.ld|l>. Details of the analysis ot the crystal field calculations for this compound will be published later. - 3.Magnetic hyperfine interaction: Fig (5) shows some representative spectra between 77K and 112K. The eight line specta obtained at the lowest temperatures suggest that the magetic field direction is not collmear with the direction of the principle axis of the electric Field gradient, and that at 77K the direction of the magnetic Field makes an angle of 60* with the princiipie axis of the electric field gradient. Fig(6) shows the variation of the magnetic Field strength H„ with temperature. This variation shows the sublattice magnetization near the critical temperature and may be represented by the equation of the form (19) M(T) - (1-T/T„)» (4) where ft is the critical point exponent and is found to be =1/2 and 0.333 and for molecular field theory and Ising model respectively.We have tested these exponer.s ; TABLE (I) Region (a) (b) Temperature rang* 78-740 K 240-320 K Center shift (CS sn mm/s) 0.727(at 77K) 1.015(at 297K) Ouadrupote splitting(mm/s) 2.54 2.198 CS regression 1:5*1.i64-4.996xio* T CS=1.1592-7.175X10"* T ln|(A(T)/A(112)) = 0.6578-5.018X10*T 0.774-2.7778X10^ Debye Temperature (K) 166±3 155 i2 Hyperfine slpmitng (kOe)at 78K 178i2*(at T * 77K) 0.0 -680- 2.60 80 120 160 200 240 280 320 Temperature I K ) Fig A Temperature variation of the quadrupole splitting between 112K -681- ...... &**>" >••••* f\*l •**•" \r ,' * I •i 1,1 '•"" •• i/! — If J: i 1 III 1 nun I 1 i 1 i • W^S« jKHpfif* "in! f "lit sVii'i' ..i f ig.5. Mossbauer Spectra ol the hyperfine structure of Fe" in HDFC below the Ne'el temperature. -6-82- X 30»- «o Fig.6. Hyperfine effective field in (kOe) as function of temperati'ra The field values are calculated from ground state splitting. -683- REFERENCES < G Chapuis. Acta Cryst. B34. ii;6 (197S) 2. W. Depemeier, Acta Cryst. 833. 3713 (19771 3 R. Kind. S. Plesko. J. Roos. Phys. Stat. Sol. (a|. 47. 233 (1977J. 4. R. Blinc. M. Burgar. a. Lozar. J. Selmger. J. Slak, V.B. Rutar. H. Arend and R. Kind. J. Chem. Phys.. 66, 278 (1977). 5 N.8. Chanh. J.R. Hausty, A Meresse. L. Richard. M Rey.lalon. J Phys Chem Solids. 50. 629 (1989). 5 N.B. Chanh. C. Hauw. A. M^resse. M. Rey-Lalon and L.J. Richard, J. Phys. Chent Solids, 46. 1413 (1985). 7 R. MoukhUsse. M. Gouzi. N.B. Chanh. Y. Haget. C. Hauw. and A. Meresse. J. Phys. Chem. Solids. 46, 167 (19S5). 8 L. Bataglia, A Bonamartinr. G. Pekni, M.R. Cramarossa, T. Manferdint. G.C. PeMacani. A. Molori, A. Saccani, F. SandrotM and M. Brrgani. Chem. HaL 4.913 (1992). 9. R.D. VWHett and E.F. Reidel Chem Phys. 8. 112 (1975). 10 K. Trch, J. Benes, W. Halg and H. Arend. Ada Cyst, B 34. 2970 (1978). 11. M.F. Mostala. MA Senary and Mil Abdet-Kader. Phya. Lett (A) 82.350 (1981). 12. M.F. Mostala. MA. Senary and M.M. Abdel-Kader. Physica. 112B. 179 (1982). 13 MA. Semary, M.M. Abdel-Kadef and M.F. Mojuta. Chem. Phys. Lett.. 96. 641 (1983). 14. A Lavsfflc C. Fife*. R. Btnc. H. Arend and R. Kind, Scfd Sole Comm.. 20,127 (1976). 15. M. Khanzi and M O. SteMty. Sold State Comm. 74.33.1 (1990). 16. J. Danon. Application ot the Mossbauer eNed in chernsby and sokl stale physics. Tech. Repr. Ser., International Atomic Energy Agency SO. 89 (1966). 17. Advances*! Mossbauer spectroscopy, B.V. Thosar, P.K Iyengar, J.X. Snrassaira and Bharasava, Etsavier Amsterdam. The Netherlands. P. 6 (1963). 18. R. Ingaaa, Phys. Rev., 135. A 787 (1964). 19 LJ. Oejongh, A.R Miedema, Advances in Physics. 23,1 (1974). -684- First International Spring School & Symposium on Advances in Materials Science (SAMS 94) 15-20 March 1994 INFRARED ABSORPTION SPECTRA OF EGYPTIAN SERPENTINE ROCK By E. M. A. KHALIL Spectroscopy Dcfrnrtmcnt, Physics Division, National Research Center, Dokki, Giza, Egypt Abstract Infrared absorption spectra of natural Egyptian serpentine rock were recorded in the frequency range 2004000 cm-1, and through the temperature range 2()0°-l000oC Tor 2 hours. The chemical bonds within the lattice structure of serpentine nrc intcrprctalcd quantitatively. From the intensities and frequencies shift or the characteristic infrared bands spectra, through heat treatment, we can detect the first cndotlicnnic reaction due to the loss of hydrosyl bonds up to 700°C. and rccryslnllization for forstritc refractory .it X(K)"C. with the characteristics infrared bands at al 506 - 521 cm-1 of Si-0 bending vibrations, at 880-888 cm-1 of Si-O-AI vibrations, and at 1000-107.') cm-1 of Si-0 si retching vibrations. The infrared bands at 350-398 cm-1 (Si-0 bend.), and 400-448 cm-l (Si-O-Mg) nrc sensitive to follow (he phase transformations through heat treatment, where Egyptian serpentine rock of chemical composition Mgf,. (Si^m) (OHR), or Mgi (S12O5) (OH)4 or 3MgO.2SiO2.2H2O magnesium hydrosilicatc, trioclahcdral 1:1 family layer lattice silicate still with its lattice structure till 700°C. cndollicrmic reaction due to the deformation ofhydroxyl bonds, and (hen formation of forcstritc of clicmic.tl composition 2 MgO. S1O2 from 800°C, and the infrared absorption Spectra of Egyptian Serpentine detect a minor quantities of AI2O3 and this phase of Serpentine is Anligoritc Mgj (Si2O«0(OH)4 trioctahcdr.il 1:1 family and high purest related to its standard. Most of the intensities of (lie characteristics infrared bauds spectra nrc decreased as temperature increases up (0 l(l(IO°C. INTRODUCTION Compact beautifully coloured varieties of serpentine arc used as facing stones, and for fashioning ornamental objects, additions of some magucsile to serpentine arc used in the production of forslcriic refractories of high melting point and good load bearing capacity. They arc characterised respectively, by volume stability, low thermal expansion, good resistance to thermal shock, and excellent cnishiug strength. In addition, these materials can be preferably used for most purposes in which basic refractories arc required. Forslcrile bricks arc used in steel industry and have also proved satisfactory for lining rotary kilns used in calcining magncsilc and dolomite. In the burning /one of cement kilns, forsicritc bricks have exhibited (lie important characteristics of taking a protective coaling of cement clinker which is a practical necessity for kiln refractory efficiency. Serpentine also finds application in chemical industry as a stock material for magnesium compounds. -685- In this paper infrared absorption spectra were recorded for fibrous antigoritc, plnty antigorite, and duysotile as different phases of serpentine (1) which have essentially the same chemical composition, but slightly different structure lowering symmetry and morphology. Ail absorption bands of antigorite could be found in chrysotilc with slight shift in positions. The clay minerals undergo on heating dehydration and dchydroxylation followed by lattice destruction and the phase transformation to mullitc (2-3). Infrared absorption spectra are used as a quantitative method in establishing the orintntion of O-H bonds in layer silicates as the hydroxyl strctclung vibratios provide an elucidation on both molecular structure and properties of clay minerals. The orientation of the O-H groups depends on the occupancy of the octahedral positions. Minerals representing (he dioctahedrai and trioctnhedral analogues of the 1:1 and 2:1 clays can be differentiated from each other mainly in the region of O-H frequencies, and their positions in these frequencies region are different The shift in frequencies of the O-H stretching vibrations is influenced by several factors, such as the preparation method of the synthetic minerals and die sources of sedimentation of natural clays, because any other mineral or minor amounts to traces of foreign ions can affect the shift of O-H stretching virbrations in one clay. The influence of particle size on the vibratioal frequencies and intensities of infrared absorption spectra were studied by several workers (4-5). It was indicated that the appearance of infrared spectra is influenced by the well- known Cluistianscn effect which is depended on particle size (6). The present authors (7- 12) study the effect of particle size and temperature of recording in the temperature range 27°-200°C, on the frequencies and intensities of AI? (Si70 The samples were prepared for infrared measurements by using the KBr disc tcclinique. The constant factors of specimens used were as follows: weight of spec-pure KBr 398 mg, weight of serpentine powder 2 mg, total weight of specimen disc 400 mg, -686- time of mixing 15 inin., pressing under vacuum (7xl()' kg/cm2) was held for 5 minutes, and then slowly released. Finally, the disc was taken Tor data recording to a Philips PU 9712 IR spectrometer, with an IRIS data station. The fully automatic double beam spectrophotometer was used in the frequency range 2(l()-4()()() cin-l. RESULTS AND DISCUSSION Effect of heat treatments on (he infrared absorption spectra of Egyptian serpentine |Mg-3(Si205)(OH)4] trioclahcdral 1:1 family arc shown in Table I and Figures la, lb, lc and Id. The infrared absorption spectra in the temperature range 200o-600°C arc the characteristic of the chemical bonds, within the serpentine lattice structure. In this range of temperatures there is no change in the 1:1 trioctahcdral lattice slniturc of serpentine. The positions of the characteristic infrared absorption bands arc not affected by firing till 600°C. except for the infrared absorption baud at 397 cni-1 shift to lower frequency 375 cm-1 as the temperature increases to (i(H)°C. The force constant is decreased for Si-O-Si bending vibrations. The characteristics of the infrared absorption bands in this temperature range arc as follows: 1. The infrared absorption band at ca. 375-397 cm-1 belong to Si-O-Si bending vibrations mode. 2. The infrared absorption band at 447 ~449cm-l is attributed for Si-O-Mg bending vibrations. 3. The infrared absorption band or R-V-OH (and ofSi-O-R^, also for Si-O-AI) at 561- 563 disappeared at 700°C, also the infrared bands at 3580 cm-1 and 3690 cm-1 arc removed at 700°C, which arc due to 0-11 stretching vibrations, and this indicates the cndollicrniic reaction due to dclvydroxylalion and deformations of the 1:1 trioctahcdral lattice structure. 4. The infrared absorption bands at 614 - 615 cni-l. 985- 988 cm-1. and 1075 cm-1 arc attributed to Si-0 linkages. 5. The absorption band at 3420 cm-l is an indication for an O-H bond which is due to both absorbed and adsorbed water of hygroscopic KBr (also this is attributed lo hydrogen bond). The absorbaucics of the infrared absorption bands chosen for quantitative studies arc decreased as the temperature increases and this mainly in a linear relationship. At 700°C a suddcen change in the spectra, due to dchydroxylnlion and a phase transformation followed by rccrystallization offorstcritc (2MgO-Si02l. This refractory material is formed in tli temperature range 800o-l0D0°C. with the characteristic chemical bonds of its network lattice structure at respectively 350-360 cm-1 (Si-O-Si bend.), 400-403 cm-1 (Si-O-Mg bend.), 506-521 cm-1 (Si-O-AI ), 609-612 cm-1 (Si-O bend.). 885 - 8X7 cm-1 (Si-O-AI). 957-958cm-l (Si-OStr.). 1003-1015 cm-1 (Si-O Sir.), and 1077-1092 cm-1 (Si-O Sir.). -687- wo—J&> iSa jsbs 71&5—S?I~55T WAVEMJMBER (Crrr') Fig. I. Infrared absorption spectra of natural Egyptian serpentine rock.in the temperature range (200°-1000°C) for 2 hours. -688- f 95J—mcrnU rr- SV—O , ( loTS-WZCm1) 6o9—61§'—0 5 Cm-J' ), Si—0 Zoo Aoo boo too looo ( 200 *» 600 600 looofoc) » Tennp«ralfe) Temperature Fig.2D. Absorbancics of Si-0 Sir. .Si-O-AI. and Si-O Fit2A Absorbancics of Si-0 bending (350-398 cm'1) bend, of Egptian Serpentine rock ihiougli heal of serpentine rock through heal treaimcnt (200*- U- .lUncnl COO'-IOOO'C) foi 2 hours. 1000'C) for 2 hours ON 1ooo-V>l5 -o-AJeBrMa, Si-O—MQvlbend o txms^psi' < ST- •e o. ritjt ?o > r i • • J. 1- V. Slubican and R. Roy, Amcr. Miner., 46, 32-51(1961). 2- I I.J. pcrcival and J. F. Duncan. Amer. Ccrant. Soc, 57 [2J, 57-6»(1974 ) 3- K. J. D. Mackenzie, Aincr. Ccrain. Soc, 55 {2|. 68-71 (1972) 4- J. M. Hunt. M. P. Wishcrcd andL. C. Bonliain. Anal. Clicm.. 22, M78- 1498(1950). 5- C. Fanner and J. D. Russell, Spectrochcmica Acta, 20, 1149-1173 (1964) 6- V. C. Fanner and J. D. Russcl. Spectrocliemica Acta. 22.389-398 (1966) 7- E. M. A. Khaliland A. A. EI-Kolali, CcraiimrgiaTcnslsgiaCcramica, VII [31.140-148(1977) 8- E. M. A. Khali! and A. A. El-Kolaii, Egypt J. pliys., 8 [l], 59-75 (1977). 9- E. M. A. Khalil and A. A. El-Kolali, Egypt J. pliys., 8 [1], 77-82 (1977). 10- E. M. A. Khalil, The First Arab Biologist Congress. Alexandria 26-30 October, 8-13,26-30 (1977). 11- A. A. El-Kolali and E. M. A. Khalil, Central Glass and Ceramic Bulletin, 25 [1J. 16-22(1978). 12- E. M. A. Khalil and M. M. K. Aouf, Egypt J. phys, 16(21.361-372 (1985). -691- First International Spring School & Symposium on Advances in Materials Science (SAMS 94) 15-20 March 1994 INFRARED ABSORPTION SPECTRA OF SILICA BY-PRODUCT OF THE EGYPTIAN FERRO-SILICON FACTORY By E.M.A.KHALIL Spectroscopy Department Physics Division National Research Center, Dokki, Giza, Egypt Abstract The infrared absorption spectra of Silica funic, which is a by-product of the Egyptian Ferro-Silicon factory are reported in the frequency range of 200 to 4000 cm-l.Thc behavior and the characteristics of the infrared absorption spectra arc recorded through heating in the range 200-I000°C for 2 hours. The results indicated that the relation between absorbencies and temperature of the infrared bands at 470-488 cm-1 (Si-O-Si bend.), at 791-806 cm-1 (Si-O-Si Str.), and at 1110-1125 cm-1 (Si-0 Str.) arc due to quartz SiO? which formed with its characteristic network structure up to 600°C. At temperatures above 800°C tridymitc is rccrystallized with its characteristic chemical bonds. From the frequency shift and intensities of the infrared bands two sudden effects due to phase transformations were delected, the first of which found as a change in intensity is due to the transformation from OC -to-p-quartz at 500°C, and the second one is due to the formation of (X-tridymite at 800°C. The intensities of the infrared bands at 470-488 cm-1, at 791-806 cm-1, and at 1110-1125 cm-1 are decreasing ?s the temperature increases, and no degradation or deformations for the network structure for these samples were found which were fumed and collected due to grinding processes from the atmosphere of the Egyptian Ferro-Silicon factory. The samples have a high purity of 95.5% SiOj. The characteristic bands of the SiOj, tetrahedra are fixed and calibra^d to that standard through heat treatment from 20()°C to 1000°C with its change transformation. The results arc similar to those obtained for the mineral quartz. The economic importance of these or silica fume materials is related to its very small grain sizes (3800 cnvVg) and. consequently, to its activity as a ceramic material. INTRODUCTION Considerable research and development efforts arc focusscd on the evaluation and utilization of silica fume which are mainly due to economical and ecological considerations. Silica fume has several potential uses in the production of fcrro-silicon materials, and these uses can be classified into those which require high temperature treatment, and those which arc produced at room temperature, as building materials. All the room temperature reactions utilize the pozzolanicity of silica funic, and this property depends on variety of factors like chemical composition, fineness, and structure of silica funic. Silica fume is a very fine siliceous powder (Blaine surface area 3800 cm-fyg), obtained as a by-product, during the production of fcrro-silicon environmental condition, and is characterized bv its small spherical particles. It is considered as a highly active -692- material (silica), due to its very high surface area and its totally amorphous character. Quartz (SiO?) appears as a scries of polymorphous modifications in crystal structure. They stand completely apart from other oxides, and are directly related to silicates concerning their crystal structures. Some authors classify this group of minerals with the silicates, since quartz is chemically a typical oxide with the three principal polymorphous modifications of SiO?, i.e., quartz, tridymitc, and cristobalitc. The scries of polymorphous transformation of S1O2 is shown as follow: a-quartz < >7?'c » 0-quartz < m'c ) p-tridymhc< MTOC >• (l-cristobalite ( l7l3'c > melt Moreover, tridymite and cristobalitc undergo cnantiotropic transformations in the region of low temperatures in the strongly supercooled state. tt-tridymit<:< "oc > p-tridymite and a-crystohalile< IM'ITOC > p-crislobalitc The present author (1-6) studied the vibrational spectra of S1O4 tetrahedra through heat treatment from 110°-1500°C for combined silica in crystalline clay minerals and also for glassy phase (7 - 20). All the natural modificaions of SiO? except opal and lechatelieritc (amorphous quartz glass) are characterised by a crystal structure, comprising of silicon-oxygen tetrahedral networks, which share comers, i.e., each silicon ion is surrounded by four oxygen ions, and each oxygen ion by two silicon ions. The structures differ in the spital orientation of the (etrahedra. The dynamical origins of the observed SiO? infrared bands at 450 cm-i, at 800 cm-1. and at 1100 cm-1 have been discussed in the light of model calculations (21 -23). This work has the aim to check silica funic as a by-product of Egyptian fcrrosilicon companies through heat treatment by infrared absorption spectra. EXPERIMENTAL The available silica fume as by-product of Egyptian fcrro-silicon factory was heated, from 100°C up to 1000"C for 2 hours, in an electric muffle furnace. After heating the samples to the desired temperature with a soaking time of two hours, the furnace was left to cool gradually. This was carried out to keep the crysiallinityofthc fired samples. The pre-fired samples were ground in an agate niortcr and sieved to grain sizes of 45 u.m. These grains were dried at 110°C for 6 hours, and then kept in a dessicator under vacuum to avoid the effect of humidity. The samples were prepared for infrared measurements, using the KBr disc technique. The experimental constant factors ofthe specimens used were as follows: Weight of spec pure KBr 398 mg. weight of silica fume powder 2 mg. total weight of the specimen disc 40O mg. time of mixing 15 minutes, and .respectively, pressing under vacuum (7x10 ' kg/cm2) was done for 5 minutes, and then slowly released. Finally, the disc was taken for data recording to a Philips PU 9712 1R -693- spectrometer with nn IRIS data station. The fully antonialic double beam infrared spectrophotometer was used in the frequency range 21)0-4000 cm-1. RESULTS AND DISCUSSION Tltc cfTcct of licat treatment for two hours, in the temperature range 200-lf)()()°C. on the characteristics infra rod absorption bands of silica funic arc shown in Tables I and 2 and Figures 1,2a, 2b, and 2c. These investigations lead to the following results: A) Phases present behveen 200° and 600°C: There arc three characteristic infrared absorption bauds in the temperature range 20O-60O°C, at 476 - 478 ciu-l orSI-O-Si bending vibrations, at Hot) - K06 cin-l orSi-O-Si stretching vibrations, and at 1120 - 1125 cm-l ofSi-0 sirclching vibrations. The positions of these bands arc not affected by firing in this temperature range, but the nbsorbaucics at 476 • 47X cni-l. and at 80(1 -806 cm-l decrease as the temperature increases. However, the intensity of the bnnd nt 1120 - 1124 cm-l was not afTcclcd by firing such that in this temperature range, (here is no change in the network structure of a- or |i-qunrtz. B) Phases present at 700°C: Two infrared absorption bands appeared at 7l)0°C namely at 1.1X5 cm-l due lo SiO., tctraltcdra with four bridging oxygen inns, and at I42X cm-l for Si04 with one non- bridging o.NNgcn ion. Al higher temperatures there is a sudden change in the >pcclr,i due to (he phase transformation of quartz to Iridymitc. C) Phases present between 800° and 1000°C: The fired silica funic, in (he temperature range X()0o-900oC. showed (he following infrared absorption bands al 313 cin-1 (Si-0bend.), at 470-474 cm-l (Si-O-Si bend). 669 cm-1 (Si-0 bend, of (ricolachcdral structure). 8(H)cm-l (Si-O-Si sir.). 873-875 cni-l (Si-OFc3+ sir). 1025-1099cm-l (Si-0 sir.), 1110-1120 cm-l (si-0 str.). 1384 cm-1 (silicon with four bridging oxygen ions) and at 1425 cm-1 (non-bridging oxygen ion). The infrared absorption band at 1620 cni-1 arc typical for H-O-H bending vibrations, due to both the adsorbed and absorbed water of hygroscopic KBr. The frequency positions of 1.184 cm-l. and 1425 cm-l may be due to hydralcd iron oxide. At l()()()°C the infrared bands al 4X6 cm-1 (Si-O-Si bend.), al 791 cm-1 (Si-O-Si str.) and at 1100 cin-l (Si-0 sir.) arc characteristic ofSiO^ Ictrahcdra in the crystalline glass stale (7 - 20) whose chemical batch composition contains Si02- CONCLUSION The intensities of the infrared absorption bands at 470 - 488 cm'1 (Si-O-Si bend.), and al 791 - 806 cm"' (Si-O-Si str.) arc decreasing as the temperature increases up lo l()00°C. showing two sudden excels al 500°C and at 8()t)°C. and of|»-qunrl/. lo tridymitc al 80()°C. In the interval 8()()°Clo I()()I)°C a stable phase, of tridymilc forms in a linear -694- Table 1: Effect of heat treatment for 2 hrs on infrared absorption spectra of Egyptian silica fume. Frequencies in cm'1 Temperature Approii. form of vibration 200C I 300t 400'C 500C 60(lt 700"C Root: 900t iooot 313 Si-0 bend. w 488 478 476 477 476 473 474 470 486 Si-O-Si bend. vs vs s vs s s s s m 669 Si-0 bend. w 806 804 80S 800 80S 800 800 800 791 Si-O-Si Str. m m • w m w w w w w 873 875 Si-0-Fe3+ Str. w sh 1025 1099 Si-0 Str. sh sh 1125 1125 1124 1125 1120 1120 1110 1120 1100 Si-0 Str. vs vs vs VS VS vs vs vs s 1385 1384 1384 S1O4 and hydrated iron oxide w w w 1428 1425 1425 Si04 and hydrated iron oxide w w w 1620 1620 1620 1620 1620 1620 | 1620 1620 H-O-H bend, of H2O abs., and ads. w w w u V w 1 w w Vs = very strong, S = strong, M = medium, W = wcak, Sh = Shoulder Table 2 : Effect of temperature on frequencies and nlisorhancics of S1O4 tctralicdra of infrared absorption bands of Egyptian silica fume. Frequencies in cm"' Temperature Approximation form of vibration 200C 300C 400C 500*C 600C 700*C 800C JOOC IOOO'C 481 478 476 477 476 473 474 470 486' Si-O-Si bend. 806 806 80S 800 •805 800 800 800 791 Si-O-Si Str. 1125 112S 1124 112S 1120 1120 1110 1120 1100 Si-0 Str. Absorbineies-logio I — V I J Approximation form of vibration Temperature 200b 300C 400C 500*C 600C 700*C 800*C ?oo'c 1000'c 3000 1.850 .904 2 420 1.143 1.244 1.785 .980 .390 Si-O-Si bend. 074 454 .286 .542 .399 421 .610 .459 .268 Si-O-Si Str. 3 000 3.000 3000 3000 3.000 3.000 3000 1841 .513 Si-0 Str. 200 1000 1500 ,2000 r-.r.,« .WWEHUMBER ( CnT«) FIG( 1) 697 A( Si-0 —- Si bend-)47o-4oo \* rl o Sbr Cm c crP. o o -MS .Pqar.tz, '.Trip'ymite ^ F'9(2aj2oo. too boo boo iooc(0C) Temperature str079l—806 CrrP Fjg,2b2oo z,oo -boo 800 looo(°c) Temperature puartz •0 • • Si—0 str.) Cm1 o o JO 1 1 1 c^/^200,400 **? 800 1oo&o % rig(2C) Temperature * rw -698- relation. The inltarcd band at 1110-1124 cm'1 (Si-Osti.) shows also a linear relation between temperatures from 2UU°C to 8()l)°C. Their absorbancies indicate that the quartz network struciuic is not affected up to 800°C. At temperatures above 800°C quartz is transformed to uidynuic where the absorbartcics decreased with a linear relationship as the temperature increases. The crystalline phases stable at auuposphcric pressure arc quartz in the temperature range up to 867°C, iridymitc from 867° (o ]470°C, and cristobalitc from 1470° to 1723°C. respectively. Rclow 172.1°C fused silica is mclastablc. Devitrification ol fused silica usually results in the formation of cnstobalitc, at all temperatures in the range from 1000°C to the melting point. Cristobalitc is not the stable pliase below 1470°C which explains that the structure of vitreous silica is crisiobalitc-likc. so that tlic formation of cristobalitc nuclei requires the least drastic rearrangement of the structural units. -699- REFERENCES 1. A.A. El-Kholaii and E.M.A. Khalil, The Ceramic Soc. of Japan, accepted for printing. November 1977). 2. E. M. A. KhaJil and A.A. El-Kolali. Cerainurgia Tcchnolcgia Coramica, \'U (3). 140-148(1977). 3. E.M.A. Khalil and A.A. El-Kolali. Egypl. J. Phys.,8(l), 59-75 (1977). .1. E.M.A. KJialil and A.A. El-Kolali. Egypt. J. Phys.. 8 (I), 77-82 (1977). 5. E.M.A. KJialil. Tlie first Arab Biologists Congress, Alexandria 26-30 October, pp. 8-13. 26-30(1977). 6 A.A. EI-KJioli and E.M.A. Khalil. Central glass and Ceramic Uiillclin. 25 (I). 16-22(1978). 7. KM. El-Badry, M.S. Afifi. F.A. Moustafa and E.M.A. KJialil. Ccniral Glass and Ceramic Bulletin. 30 (1). 11-17 (1983). 8. K.M. El-Dadry. M.S. Afifi. F.A. Mouslafa and E.M.A. Khalil. Egypt. J. Phys.. 15(1), 143-52(1984). 9. K.M. El-Badry. M.S. Afifi. F.A. Moustafa. and E.M.A. Klialil. Ccniral Glass and Ceramic Bulletin, 31 (1-4), 23-27 (1984). 10. F.A. Moustafa. K.M. El-Badry, M.S. Afifi and E.M.A. KJialil. Accepted Indian J. of Pure and Applied Physics (1985). 11. E.M.A. Khalil and M.M.K. Aouf, Egypt. J. Phys.. 16(1), 135-142 (1985). 12. E.M.A. Khalil and M.M.K. Aouf, Egypt. J. Phys. 16 (2). 373-378 (1985). 13. E.M.A. KJialil and M.M.K. Aouf. Egypt. J. Phys.." 16 (2), 361-372 (1985. 14. E.M.A. Khalil. and M.M.K. Aouf, Egypl. J. Phys.. 17 (1). 255-261 (1986). 15. E.M.A. Khalil and M.M.K. Aouf, Egypt J. Phys.. 17(1). 249-254 (1986). 16. E.M.A. Khaiil. Bull. NRC Egypt, 17 (3), 181-195 (1992). 17. E.M.A. Khali!, 4di Nat. Phys. Conf., Cairo University, 28-30 November, pp. 157-166(1992). 18. E.M.A. KhaJil land M.M.K. Aouf, 4th Nat, Phys. Conf.. Cairo University. 28-30 November, pp. 167-176 (1992). 19. E.M.A. Khalil and M.M.K. Aouf, 4th Nat. Phys. Conf. Cairo University. 28-30 November, pp. 177-186 (1992). 20. E.M.A. Khalil and M.M. Aouf, 4th Nat. Phys. Conf.. Cairo University, 28- 30 November, pp. 187-196 (1992). 21. R.J. Bell, N.F. Bird and P. Dccan, Phys. Chem., I, 299 (1968). 22. K. Nakamoto, Infrared Spectra of Inorganic and Coordination Compounds. John Wiley and Sons, Inc., New Yoik (1963). 23. J.M Stcvcls. J. Soc. Glass Techno!.. 35.284-288 (1951). -700- First International Spring School & Symposium on Advances in Materials Science (SAMS 94) 15-20 March 1994 RELATION BETWEEN MAXIMUM DEGREE OF HYSTERESIS AND BOTH TEMPERATURE AND INITIAL PRESSURE DURING HYDRIDING OF LaNi'5 S.L. ISAACK*, E.A. KARAKISII**, H.I. SHAABAN*, •ATOMIC ENERGY AUTHORITY,ALEX.,EGYPT **IRGS,ALEX. UNIV..EGYPT Abstract: In the LaNi5-H2 system constant volume hydriding / dchydriding cycles were performed in the initial pressure and temperature ranges from 1.1 to 4.2 MPa and from room temperature up to 523 K respectively. It was observed that the maximum degree of hysteresis decreases with temperature and pressure increases. Furthermore, it was found that during the activation present at constant temperature the degree of hysteresis decreases by cycling at low temperatures, but the phenomena was found to be reversed at 373 K. The degree of hysteresis was also found to decrease cyclcwisc when a hydrogen inter cycle thermal evacuation step was carried out (open cycles). However, the degree of hysteresis remained more or less unaffected after repeated hydriding/ dchydriding cycles were achieved without performing the hydrogen evacuation step (closed cycles). Hence, it was concluded that the internal stresses build up is higher during closed cycling rather than during open cycling whereby stress relief takes place by thermal treatment during evacuation. -701- UNTROOUCTION Hysteresis is one of the phenomena associated with the hydrogen storage process in bolh metals and intcrmctallic compounds . The phenomena has the drawback of decreasing the efficiency in practical applications of these hydrides. Hysteresis is usually described as a pressure difference between the equilibrium pressures of hydride formation and decomposition . The phenomena was reported by many authors as early as 1925 [I] for the palladium • hydrogen system. Kiijpcrs and Van Mai [2] investigating the LaNi5-H2 system found - although unremarkable at low pressure and temperature - that the extent of hysteresis increases with both temperature and plateau pressure . Birnbaum ct al [3] attributed this bolivian to the generation of dislocations along each plateau pressure and their suggestion was evidenced by the electron microscopic investigation carried out by Wise et al [4] on the Pd\H system. Flanagan ct al [5] calculated the enthalpy loss which associates the free energy loss in a hydriding/dchydriding cycle and found that the difference between the cooling and heating press corresponds to twice the dislocation enthalpy . The same authors [6] studied also the effect of annealing on plateau hydrogen pressures and hysteresis in LaNi5-H system and found that the absorption plateau pressure is largely affected by the annealing treatment of the active LaNi5. Flanagan and Clcwly [7] presented also a detailed model for hysteresis based on the enthalpies requirement for dislocation production during both hydride formation and hydride decomposition. Wagencr [8] showed by statistical thermodynamics for the Pd-H system that the energy loss in the hysteresis loop, RT In (Pf/Pd) .should increase markedly with a temperature decrease . Wcifang et al [9] determined the energy stored in activated LaNis and found that the hydrogen pressure needed for hydride formation is greater for annealed LaNi5 than for the activated material. Recently, Hyo- Jun et al [10] found that after 3150 hydriding/dchydriding cycles the plateau region in the P-C isotherm curves for the LaNi5 -H2 system disappeared completely and its slope increased sharply with a decrease of about 87% of the hydrogen storage capacity. All the above studies cited above were performed under isobaric conditions .Since -702- the technical application of hydrides makes if necessary (o go through a constant volume process the authors described in a necessary (111 the present hysteresis loops when (he LaNi5 was activated isochorically. The present paper describes the cfTcct of both temperature and initial pressure on (he shape of the developed hysteresis loops in this system The intcrmctallic compound LaNi5 was selected for this investigation because of. respectively data availability, and general interest. 2. EXPERIMENTAL Two sets ofLaNi5 samples (2.5 g each) were introduced in a standard stainless steel Sivcrt's apparatus and were subjected to the following pretrcatment activation procedures: a) The first set of samples was subjected to repeated hydriding dchydriding experiment by only cooling alone slowly from the designed initial temperature down to room temperature , then reheating to the starting temperature without withdrawing or introducing any hydrogen from the capsule (closed system). b) The second set of samples was also hydrided by cooling from the required initial temperature then the formed hydride was thermally dehydrided by reheating (o the starting temperature then evacuation at 10 Torr at this final temperature (open cycle) . A fresh hydrogen aliquot was introduced in (he subsequent hydriding process. The developed hysteresis loops during activation by the two methods as also the hysteresis loops for active compounds and (he degree of hysteresis were plotted against temperature using computer software for best fit. The degree of hysteresis was calculated from the equation (Degree of hysteresis = 1/2 RT In Pf/Pd) where Pf and Pd denote hydride formation and decomposition pressures at constant temperature respectively, R is the international gas constant and T denotes the absolute temperature. -703- 3. RESULTS AND DISCUSSION 3.1. Hysteresis during activation of LaNij Figure 1 shows the hysteresis loops developed during activation of the LaNif samples using closed cycle method in hydrogen of initial pressure of 3.4 MPa and at 523K . It could be seen that there is a cyclcwisc loss of initial pressure which denotes that an appreciable amount of hydrogen corresponding to this pressure loss is lightly bound to the compound and could not be driven off thermally. In other words an increasing amount of thermally stable hydride is formed cyclcwisc. The areas of hysteresis loops arc found to be more or less identical in spile of the cyclcwisc decrease in initial pressure . Closed activation cycles at higher pressure (4.2 MPa) and at higher lempciature (623 K) resulted in a pronounced cyclcwisc loss in hydrogen storing capacity . Contrary, open activation cycles (Figure 2) resulted in an active LaNi5 compound keeping almost its original hydrogen storing capacity . This leads to concludc'that thermal dchydriding under vacuum or low pressure is a beneficial process since it reserves the original capacity of the LaNis hydrides. Figures 3 and 4 show the plots of the degree of hysteresis as a function of temperature. It could be clearly seen that the maximum degree of hysteresis occurs at 370 K for samples activated in open process and at 390 K for samples activated in a closed cycle. Also, the degree of hysteresis increases cyclcwisc during aclivation.The increase is more pronounced with samples activated in closed cycles. The observed increase in the degree of hysteresis in the open cycle is only apparent, but this is not trucly the case when considering the simultaneous decrease in both the nominator (Pf) and denominator (Pd ) in the mathematical term. As the degree of hysteresis could be considered as a measure for dislocations in the crystal lattice [3,7] , these results may presume that the maximum degree of dislocations during activating the LaNi5 take place around these temperatures .The greater degree of dislocations happening in the LaNi5 when activated in repeated closed cycles (when compared numerically with value of open cycles) denotes the cotinous presence of hydrogen at higher pressure around and in the lattice a matter which enhanced the -704- J. 2 £™ In Cycle No. ceoool a. ••••• | 3 A44A43 I » J.l.l J-UUI-L.tl.IJ ,1.1.1 II ,1.1.1 I I I i-1 I IJL.I.JJ.JJ IL.U t III. esu Temperature (K) Fig .1. Hysteresis loops developed during activation ofLaNi5 using closed cycle in hydrogen oT initial pressure of 3.4 MPa and at 523 K ».» a a. n V I. o. «i _3 "S .2 *«» t • i t I i « • t . i i • • I i till i*,7V. I J50 «50 5V> Tetimeralure (10. Fig 2. Hysteresis loops developed during activalion of LaNi5 using open cycle in hydrogen of initial pressure of 3.4 MPa and at 523 K -705- i. in •JH r.iKHi 4) u«» Km (1 to oV 7tK)ll 0 0 .1-ULLl-tXX.tljXlJLl S50 Toni|)cralure(K) Fig.3. Varialion of Hie degree of hysteresis wild Icpcralnrc or LaNi5 samples acUvalcd using closed cycle in hydrogen with nn initial pressure of 3.4 Mpa. Tt!tu|>cmt«irc(K) Fig. 4. Variation of the degree of hysteresis with temperature of LaNi 5 samples activated using open cycle in hydrogen of initial pressure of V4 MPa -7l)r»- dislocation al higher temperatures. Al temperatures higher than 450 K Ihc degree or hysteresis was found to decrease cyclcwisc Samples activated using the open method showed a cyclcwisc decrease in the degree of hysteresis over the whole range of temperature tested. 3.2. Hysteresis in active LaNi^ -H2 system As it was found experimentally that Ihc closed cycle method has a detrimental effect on Ihc hydrogen storage capacity of the l.aNij, only samples activated using the open cycle method were used to study the effect of (ho initial hydrogen pressure on the variation of the hysteresis with temperature in the LaNi5 -H2 system Figure S shows the obtained hysteresis loops for active samples when the initial pressure was varied from I.I MPa to 3.4 MPa. Figure 6 shows the variation in the degree of hysteresis with temperature in the range of 320- 523 K . It could be clearly observed that the maximum degree of hysteresis shifts towards higher temperature with pressure decrease . A fact which denotes that the dislocations accompanying the hvdriding process must be aided by means of thermal energy when low pressure hydrogen has to react with the LaNi5 intcrmctallic compounds. This would lead to the concision that dislocations formation within the crystal lattice may be an associated phenomena in achieving the introduction of the hydrogen atoms into (he interstices of thcLaNis lattice. 4. CONCLUSION Experiments carried out in the present study pertaining the LaNi5 -H2 system lead to the following conclusions: 1. It is belter to activate the LaNi5 using the lowest possible hydrogen pressure and temperature and to dchvdridc the obtained hydride in vacuum in order to reserve its storing capacity for a larger number of cycles. -707- C/cl« Ho. . ••••• I Cycle Ho ( • •••• i =, y^y C'« rs5i= s" o«1 =/-.- •%>* •Ms Tfin|>«r>(iii e (K| Tein|ieri(ur*(K) (bi Fig .5. Hysteresis loops Tor active l..iNJ5 siiiuplcs liydridcd in hydrogen or initial pressures: a) 1-3.4 MPa , 2-2.8 MPa and 3-2.3 MPa b) 2.2-1.6 MPa and 2-U MPa Cycle Wo. 3.«IMPa*££* ' QQQQD2 Z.fJMPaQfiOflO I AA&fltA I 1-OMP.JJjgJ J 1,1 AAAAA2 500 Tciuperalure(K) Fig .6. Viiriation of the degree of hysteresis with temperature for LaNJ5 active snniplcs when liydridcd in hydrogen having different initial pressures -7(>S- 2. Maximum degree of hysteresis (dislocations within ihc lailicc) is obtained at 390 K when LdNis, is activated using the closed cycle ntclliod and at 370 K if open cycle procedure is followed. 3. LnNi5 activated by the open cycle method showed cyclcwisc decrease in Ihc maximum degree of hvtcrcsis along llic tested Icinpcralitrc range (300 -523 ). 4. The maximum degree of hysteresis decreases with pressure decrease while shifting towards higher icmpcraliircs. 5. Dislocations formation within the LaNis. crystal lattice may be regarded as an associated phenomena in achieving the introduction of hydrogen atoms into the interstices of the lattice . It may be reduced by decreasing Ihc hydriding pressure REFERENCES 1. N. A. Schollus and W.Keith Hall. J.Chcm. Physics vol 39, No.4,1963. 2. F.A. Kujpcrs and H.H. van Mai. J.Lcss- Common Mclals, 23, pp 395-398 (1971). 3. H.K. Birnbaum, M.K. Grossbcck and A. Amano. J.Lcss-Conimon Metals, 49,p 357(1976). 4. M. Wise, J. Farr, I.R. Harris and J.R. Hist, L'Hydrogcn dans les Mctcaux" Editions Science and Ind„ Paris, vol 1, PI (1972). 5. Ted B. Flanagan, B.S. Bowcrman and G.E.Bichl. Scripia Metallurgies vol 4, pp 443-447(1980). 6. Ted B.Flanagan and B.E. Bichl . J.Lcss-Common Mclals 82,pp 385- 389 (1981). 7. Ted B.Flanagan and J.D. Clcwly. J.Lcss -Common Mclals, 83,pp 127-141, (1982). 8. C.Wagner. Z.Phys. Chcm. (Leipzig), 193, P38fi (1944). -7tW- 9. Wcifang LuoandTcd B. Flanagan.J.Lcss- Common Mclals, 142, 281- 288 (1988). 10. Hyo-Jun Ahn Jai-Yoting Lcc .liu. J.Mydrogcn Energy vol 16. No. 2. pp 93 - 99. (1991). 11. S.L Isaack. E.A. Karakish . H.I.Shaaban and S.J.Kandil, Proc. World Rcinucwabic Energy Congress, Reading U.K.. 13-18 Sept- 1992 -710- £Gf5oyfS£f First International Spring School & Symposium on Advances in Materials Science (SAMS 94) 15-20 March 1994 PHASE TRANSITIONS DURING HYDRIDING ACTIVATION OF LaNi5 INTERMETALLIC COMPOUND ISAACK S.L.* KARAKISII E.A.** AND SHAABAN H.I.* *ATOMIC ENERGY AUTHORITY, ** IRGS, ALEXANDRIA UNIVERSITY, EGYPT ABSTRACT: This paper presents a trial to detect the phase transitions during the different activation cycles of LaNi5 intermctallic compound when subjected hydrogenation at high temperatures and pressures. The phase transitions were followed by using the X-ray diffraction technique. It was observed that the hydrogenatcd samples showed the presence of LaNi5 hydride plus the formation of one or more new phases. The dchydrogcnatcd samples showed the presence of LaNi5 plus the LaNi5 hydride phase in conjunction with some new reflections at various diffraction angles. The absence of oxide which might be formed during hydriding was checked by subjecting the following samples to X-ray diffraction analysis : the as received, hydrided and dehydrided LaNi5 samples treated in air at high temperature (1173KK) for 2.5 hours. The as conncrcial received sample, the differently hydrided and dehydrided samples gave only oxide phases . It could be concluded that the newly formed phases during hydriding are not an oxide or oxides of La and/or Ni. These foundlings were also consistent with the data obtained by TG and DTG techniques -711- 1-INTRODUCTION The intrinsic degradation behavior of LaNiS results in decreasing the hydrogen storage capacity of the compound . This phenomenon has been investigated by many authors [I-10] in order to clarify its causes. The different intcrmclallic compounds that arc formed by alloying La and Ni metals together with their crystalline structures were given by Bushow and Van Mai [ll|. Lynch and Rcilly (12] investigated the hydrided and dehydrided samples by X-ray diffraction and found no new phases to form upon repealed cycling at low hydride composition. Pcrchcron ct al 113] using neutron and X-ray diffraction profile analyses found (hat the LaNiS hydride gave line broadening reflections due (o (he presence of micro strains. Goodcll [7] formulated the free energy for the alloy disproportionation as: 1/3 LaNiS + H2 = 1/3 LaNiS H6 G = 7400 + 26.0T LaNi5+H2 = LaH2 + 5Ni G= 19180 + 35.4T This author and he suggested also the formation of the stoichiometric compounds La2Ni7 and LaNi which were previously investigated by Macland ct al [14] and were found to form stable hydrides. He suggested that the degradation resulted as a consequence of atoms reordering in which lanthanum and nickel atoms exchange positions in the crystal lattice. Park and lee [9] proposed that LaNi5 undergoes degradation during hydriding as a result of a combination of microphasc separation at surface and reordering in the bulk. Hyo - Jun and Lee [10] using X-ray, differential scanning calorimeter and gas chromatography techniques were able to detect Ni-clustcrs and La-Ni- hydrides in the degraded LaNi5 aflcr 3150 hydriding / dchydriding cycles. In the present work an accelerated activation process at high temperature is used in order to throw more light ondetails of in the hydriding /dchydriding process. The addition, understanding is obtained concerning the activation process and the reasons for decrease of the hydrogen storage capacity. 2- EXPERIMENTAL Stoichiometric LaNiS samples supplied by the MPD Co (England) were used in (his s(udy. A stainless steel standard Sicvcrt's apparatus was used in carrying out the -712- hydriding / dchydriding cycles. The LaNi5 siimplcs (2-2.5 g each) were subjected to hydriding in hydrogen gas having an initial pressure ranging from II to4.2MPa and by cooling / reheating in the temperature range : room temperature -1I73K . The as - received hydrided and dcliydridcd (activated at different stages) samples as also the corresponding air healed samples for 2.5 hours at 1173 Ka) were analyzed using X-ray diffraction (CU Ka) and thcrinograviuictric methods. 3. RESULTS AND DISCUSSION Figure 1 shows the X-ray diffraction patterns obtained from: a) The as received LaNi5 sample . b) LaNiS subjected to 5 cycles of hydriding by cooling from 1173K to room temperature in hydrogen of initial pressure of 3.4 PMa, then dchydriding by reheating to the starling temperature without hydrogen withdrawal. c) LaNiS sample subjected to 5 cycles of hydriding by cooling from 1173 K to room temperature in hydrogen of initial prcssurcof 3.4 MPa then dchydriding by reheating. Them annealing under vacuum of 10-3 Torr at the starting temperature for one hour. d) LaNiS (as received) sample hydrided once by cooling from 1173 K to room temperature in hydrogen of initial pressure of 4.2 PMa, then dehydrided by reheating under vacuum of 10-3 Torr to the starting temperature. It could be clearly observed that the hydrided and dehydrided LaNiS samples (b) and (c) show identical patterns with newly generated peaks not found in the original as received sample which denote the formation of new phases under the effect of the drastic hydriding / dchydriding reactions. These peaks were matched with that given by Jcong and lee [IS] and Tas and Hewu [16] fori the oxidized LaNi5 products and that given by HYO-Jun and Lee [10]. Peaks of the LaNi5 H3.6. LaH3, LaNi5 and Ni were identified. New peaks appearing at refraction angles between 10-20 and 4S-SS (2 thcta degrees) characterized the LaNi5 compound which was hydrided once at 4.2 Mpa hydrogen pressure then dehydrided. More analysis was needed in order to -71.V identify these intermediate tempcrorary peaks winch disappeared in the subscqucnthydriding / dcliydriding cycles. The X-ray diffraction patients for (a) the LaNiS hydrided for once al 4.2 MPa hydrogen, (b) part of (a) affected to dchydriding in vacuum and (c) sample (a) oxidized in air nt 1173 K were given in figure 2. From this figure it could be seen that this intermediate phase is not an oxide and that the dehydrided sample showed new peaks which were not present in the parent hydride. The intensity of the new peaks given by the hydride arc not as much as large as that given by the fully activated samples shown in figure I . Hence the activation process of the LnNiS may be due to the formation of new active phases or centers which promotes the hydriding reaction. These centers may cither be Ni clusters whose reflections appear in the activated LaNiS samples or new phases formed from the low pressure hydridablc - inlcrmclallics La2Ni7 and LaNi3 [7.12]. In all hydrided - dehydrided samples the peaks denoting the hydride LaNiS H3.6 were identified, this means that the residual hydrogen in the LaNiS is most probably in the form of this compound. Figure 3 shows (he thcrmogravinictic analysis curves for the hydrided and dehydrided LaNiS samples when carried out in static air. The curves show reactions associated with weight increase at 598 K and 853 K for the hydride sample and at 751 K and 915 K for the dcliydridcd one. These reactions arc mainly oxidation reactions leading to the La and Ni oxides as was kind by chemical analysis. The lower temperatures at which these reactions happen in the case of hydride sample •nay be explained by the evolution of hydrogen and its ignition, a matter which locally raises the temperature of the sample and bring it to the oxidation reaction temperature. These curves concluded that oxidation of these samples starts al 598 K and ends at 915 K. Figure 4 shows the X-ray diffraction patterns of the as received, hydrided and dchydrided samples after being subjected to air and heating at 1173 K for 2.5 hours to ensure complete oxidation. All the reflection angles arc identical but different from those obtained by the new peaks appearing in the hydrided and dehydrided samples (figure 1 and 2). This proves that new peaks obtained by these samples arc indicative of new phases to be investigated since some of them were not reported up to now. -714- 'IVro 'II I'll n Figure (I): X-ray diffraction patients fur LiiNi^ samples: i) As received (b) Aclivalcd through 5 cycles by cooling from 1173K in hydrogen or 3.4 MP;i initial pressure Mowed by Ihcrmaly dchydridiiig for 5 cycles, (c) Activated llirotigli 5 cycles as in (b) wilh intercycle hydrogen evacuation at II73K. d) Ilydridcd nnd dchydridcd once by cooling from 1173 K to room temperature in hydrogen of-I.2M Pn followed by vacuum dchydridiiig . -715- %Jl^%%/W WW iw^S'hk^m!>h*\ ?w%4WW^ mwm0i \mm W#w w^i^JWw^ '^^^^ —I Co ftO -IO 30 SO Two IIIMIII Figure (2): X-ray diffraction patterns for LaN'15 samples: a) As received (b) llydridcd by cooling from II73K lo room temperature in hydrogen of 4.2 MPa initial pressure (c) Sample (b) dehydrided at II73K in vacuum (d) llydridcd, then air healed nl I I7.1K for 2.5 hours. -71(v (») 59BK 603R | 1 1 1 1 1 1 1 1 290 373 «73 573 673 773 073 973 1023 Teui|)ci-oUii-e(K) (I-) 75IK OTC ~~I « 1 1 1 1 1 1 1 298 373 «73 573 073 773 B73 973 1023 Teinpcralurc(K) Fifurc (3): TG and I1TC <>r Lanij Samples (In sialic air) (a) livdridcd (b) dcliydridcd -717- I \Jww^iJV W W Wrffew^^^ Vw Ir^M^ ** V w^*w%^w;, % J WAtyWvA\^^^ GO 0l» 40 3l> ZO IO 1V»o Uinta Figure (4): X-ray diffraction patterns Tor air licalcil samples (b) As received LaNiS (b) Hydride : LiNiS 115 8 (d) Snmplc (c) dehydrided. 4. CONCLUSION The stepwise follow up using X-ray diffraction technique of the l.aNi5 inki metallic componiul samples during the activation process lo IK- susceptible lo liydridini' led to the following results: 1. An intermediate phase may form after llie first activation cycle which disappears in (lie subsequent cycles. This phase needs more investigations for identification 2. Active samples showed Ihc reflection patterns of the metals and hydrides i.e. Ni. LaH2 and LaNi H5.6 even in the vacuum dehydrided samplcs.Thc presence of these phases may be responsible for Ihc activity of LaNi5 towards the hydriding process. REFER EN CES 1. H C. Sicgcmann . L-Sclil.ipb.ich and C.r Bmndlc. Phys Rev Lett 4(1. «J72 (1978). 2. R.L.Cohcn, K.W. West and J.H. Wcrnick. J-Lcss-Common Metals 70.229 (1980). 3. R.L. Cohen. K.W. West and J.H. Wernicke. J-Lcss- Common Metals. 73 . 273 (1980) 4. R.Rummcl, R.L. Cohen, P. Gutlich and K.W. West. Appl. Phys. Lett. 40.477 (1982). 5. R.L. Cohen. R.C. Sherwood and K.W. West. Appl. Phys. Lett, 41.999 (1982). 6. R.L. Cohen and K.W. West. J. Less-Common Metals. 95.17 (1983). 7. P.D. Goodcll. J. Less -Common Metals 99.1 (1984). -719- X M.J. Bcnliam, D.K. Ross. C. Lariiguc :md A.Pcrclicron - Gucgan. Zcil. Plivs. Chcm. N.F. 147.219(19X6).' 9. Jong - Man Park and Jiii - Young Lcc. Malcr. Res. Bull. 22 455. (1987). 10. Hvo-Jun AHN and Jai-Yoiiiig Lcc. Int. J.Hvdrogcii Energy. \oi I6.NO2 . 9.1-99 (1991). 11. K.H.J. Buscliou and H.H. Van Mai. J. Less - Common Metals . 29.203-2II) (1972). 12. J.F. Lynch and J.J. Reilly. J. Lcss-Comiiioii Metals, 84 . 225- 236 (I W; 13. A. Pcrclicron- Gucgan. C.Lartiguc. J.C. Acliard, P Gcrmi and, F. Tassel-. J. Lcss-Coniiiion Metals 74, 1-12 (1980). 14. A.J. Macland, A.F. Anderson and K. Vidcin, J.Lcss • Common Metals, 45, 347- 350 (1976). 15. Jcong in Han and Jai - Young Lcc. J.Lcss-Comnion Metals. 152,329-338(1989). 16. H.Taoand Y.Hcwn. J.Lcss - Common Metals, 153, 253-257 (1989). -720- AUTHOR INDEX OF VOL. II -721- First International Spring School & Symposium on Advances in Material Science (SAMS 94) Cairo, 15 • 20 March 1994 AUTHOR INDEX OF VOL. II Abbas A. 329 Abd El-Azim M.E 351 AbdElKaderMM. 675 Abd El-Razik LD. 77, 181, 389 Abd-Elhady MJL 255 AbdelRaoufM.A 599 Abdel Reheim N.A. 265 AbdulMajid S. 667 Abu-HelalH. .651 Abulfaraj W.H. 667 Afifi H.H. 431,453, 463,483 Afifl Y.K. - 181,389 Agamy S. 181 Ahmed F.S. 45 Ahmed H.A. ,93, 113, 265 Ahmed MS. .— 609 Al'Mohamed A 559 •723- Arafat S.S. 675 Ashoub N. 643 AshourA 441, 497 Alalia S.R. 623, 633 AttiaAM 25 Attia G.A. 623 BishayAS. 77 Buchkremer H.P. 131,157 Darwish S. 463 ElAslabi A.M. 77 El-Desouky A.R. 25 El-Eskandarany M.S. 93,113 El-Essaui M. 675 El-Fefey M.E. 285 ElGhazalySA. • 363 El-Hakim SJ\. 483 ElHoute S. 297, 507, 535 El-KadryN. - 441 ElMashri S.M. 173 El-Masry M.A.A. • 517 -724- El-Mast yNA. 373 El-Nuggar M.M. 623, 633 El-Naklmli A.E 329 El-Ra/fhy S.M. 217,239 El-RafeyF.A. 265 El-Sayed AA. 61, 217, 239 El-Sayed Ali M. 507 El-Sayed M.M. 285 El-Sayed T.A 623, 633 El-Shanshoury A.l. 569,585 El-Shanshoury LA. 569, 585 El-Sharkawy AA. 623, 633 El-Shokrofi K. 25 Enrick R. 675 Fayek M.K. - 651 Gad M.MA. 217, 239 Gadalla AM. 517 Gamal G.A. 633 Ghoneim M.M. 15,61 flamnmdF.H. HI, 157.297.351 -; 2S. Hansen A' 297 HosnyA.Y. 285 Ibrahim S. 329 haack S.L. 701, 711 Kaiser M. 651 Karakish E.A. 701,711 Kassem ATA. 255 KcnawyM. 623 Klialil E.M.A. 685,691 Khalil M.S. 431 Khedr S. 9 Kliedre C. 25 Kleist G. 77 Krawczynski S. 181 Lorenz W.J. 599 Mahmoud S.A. 441,497 Mina N.K. 623,633 Mohamed K.E. 131, 157 Momtaz R.S. 431, 453, 463 •Moray S.M. - • 9, 285 -726- Mostafa M.F. 651, 675 NagatA.T. 633 Nasr eldin A.M. 61 NasrH. 363 Nasser S.A. 483 Nickel H. 77 Orabi G.I. 569,585 PachurD. 61 Ramadan R.M. 329, 339 Reutov V.F. 549 Samir Abdel-Aziin M. 9 Sayed Ahmed F.M. 651 Scherpke G. 643 ShaabanH.1. 701, 711 Smeltzer W.W. 265 Soliman H.M. 417 Sorensen O.T. 507 Soukieh M. 559 Stover D. 131» 157 TahaA.S. 297»313 Terra F.S. 453 Turkustani O. 667 WaheedA.F. 417 Zaghloul Hala S. 5" -727-