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The development, growth and oxidation resistance of boron- and germanium-doped silicide diffusion coatings by fluoride-activated pack cementation. (Volumes I and II)
Cockeram, Brian Vcm, Ph.D. The Ohio Sttte Univenity, 1994
UMI 300 N. Zeeb Rd. Ann Arbor, MI 48106
THE DEVELOPMENT, GROWTH AND OXIDATION RESISTANCE
OF BORON- AND GERMANIUM-DOPED 8ILICIDE DIFFUSION COATINGS BY FLUORIDE-ACTIVATED PACK CEMENTATION VOLUME I
DISSERTATION
Presented in Partial Fulfillment of the Requirements for the Degree Doctor of Philosophy in the Graduate School of The Ohio State University
by
Brian vern cookeram, B.S., M.8
The Ohio State University
1994
Dissertation Committee: Approved by R.A. Rapp K.H. Sandhage Advisor f (r W. Soboyejo
Program in Department of Materials Science and Engineering To my wife Teresa and son Andrew, with love
ii ACKNOWLEDGEMENT
I want to express my deep gratitude for the advice and support of my advisor, Professor Robert A. Rapp. His constant encouragement, expert advice, and critical comments on my work have been a blessing over the past three years. I thank him for the opportunity to present my work at a number of meetings and to publish in journals. It has been a real pleasure to work with such a distinguished, bright and creative scientist. I thank Madan Mendirrata (UES, Inc.) and Tom Broderick (Wright Laboratory, Wright-Patterson AFB) for providing financial support for this work. Thanks again to Tom Broderick and Jon Schaffear (GE Aircraft Engines) for providing the Ti-22Al-27Nb and Ti-20Al-22Nb alloys used in this work. Paul Bania (TIMET) donated the commercially pure titanium that was used throughout much of this work. I am especially grateful to Bill Brindley and Jim Smialek of NASA
Lewis Research Center who gave excellent advice on how to evaluate my coatings, with Ti-22Al-23Nb and monolithic and composite Ti-24Al-llNb for coating.
iii A special thanks is given to the members of my dissertation committee, Drs. Ken Sandhage and Wole Soboyejo for their advice with this work. Their hard work and insightful comments are reflected in this document. Thanks to Bud Farrar, Gary Dodge, and Henk Colijn for providing help or assistance in times of need. I want to thank my friends and fellow graduate students in the Department of Materials Science and Engineering, who have made the past three years at Ohio State a pleasant experience. I thank the members of Dr. Rapp's group, past and present: Bob Bianco, Mark Harper, Ge Wang, Patrick Choquet, Anna strawbridge, Rengan Krishnakumar, Xuejin Zheng, Ted Cruise, Chris McVay, Mulong Yu, Ms. Wu, Drs. Shen and Zhang for their friendship, help and advice during the course of this work. I am especially grateful for the opportunity to collaborate with Dr. Ge Wang, a bright mind for the discussion of ideas or theories, as demonstrated by our invention disclosure for preventing the posting of MoSi2 . Last but certainly not least, 1 thank my wife and child for enduring my frequent and long absences during this work. I also thank my parents, in-laws, and extended family for supporting roe in this endeavor. In closing, I thank God for the abilities I have been given and the opportunities provided to use them. To God be the glory, for great things he has done. iv VITA
February 24, 1967 Born-Dayton, Ohio
March 1990 B.S. Ceramic Engineering The Ohio State University Columbus, Ohio
December 1991 M.S. Materials Science Wright state University Dayton, Ohio
December 1991-present Graduate Research Assistant Department of Materials Science & Engineering The Ohio State University Columbus, Ohio
PUBLICATIONS *B. Cockeram and R.A. Rapp, in Proceedings of the 3rd International Conference on High-Temperature Intermetallies, May 15-19, 1994, San Diego. *M. Thirukkonda, B. Cockeram, M. Sagib, L.E. Matson, R. Srinivasan and I. Weiss, Scripta Metallurgica et Materialia*, 27 (1992) 711-716. *B. Cockeram, A.G. Jackson, R.E. Oralor, R. Srinivasan and I. Weiss, Journal of Electron Microscopy Technique, 22 (1992) 298-300.
*B. Cockeram, R. Srinivasan and I. Weiss, Scripta Metallurgica et Materialia, 26 (1992) 755-760. *B. Cockeram, M. Saqib, R. Srinivasan and I. Weiss, Scripta Metallurgica et Materialia, 26 (1992) 749-754.
*B. Cockeram, H.A. Lipsitt, R. Srinivasan and I. Weiss, Scripta Metallurgica et Materialia, 25 (1991) 2109-2114. *B. Cockeram, M. Sagib, R. Omlor, R. Srinivasan, L.E. Matson and I. Weiss, Scripta Metallurgica et Materialia, 25 (1991) 393-398.
v FIELDS OF STUDY Major Field: Materials Science
Studies in:
Chemical Metallurgy: R.A. Rapp, G.R. St.Pierre, . Akbar Physical Metallurgy: D.A. Dregia, H.L. Fraser
Mechanical Metallurgy: G.S. Daehn
Corrosion: R.A. Rapp, S. Smialowska Electron Microscopy: W.A.T. Clark, H.L. Fraser Solidification: C.E. Mobley Thermal Analysis: P.K. Gallagher Solid-State Physics: R. Sooryakumar
vi TABLE 07 CONTENTS
ACKNOWLEDGEMENTS . iii
VITA ...... v LIST OF TABLES . . ix
LIST OF FIGURES . . XV
ABSTRACT XXV CHAPTER PAGE I. INTRODUCTION ...... 1 1.1 Requirements of Aerospace Materials 1 1.2 Titanium-Base Alloys ...... 2 1.3 Protection of Niobium Alloys . . . , 6 1.4 Pack Cementation Coatings ...... 7
II. LITERATURE SURVEY ...... 12 2.1 Halide-Activated Pack Cementation . . 12 2.1.1 Pack Siliconizing ...... 12 2.1.2 Pack Boriding ...... 15 2.1.3 Pack Chromizing and Aluminizing 15 2.1.4 Codeposition in a Single Processing S t e p ...... 17 2.2 Thermodynamics and Kinetics of the Pack Cementation Process ...... 19 2.2.1 Single Element Deposition . . . 19 2 .2.2 Codeposition ...... 28 2.2.3 Growth Kinetics ...... 36 2.2.4 Growth of Multi-layered Coatings 40 2.3 Performance of Coatings ...... 44 2.3.1 Interdiffusion ...... 46 2.3.2 Isothermal and Cyclic Oxidation 47 2.3.3 Substrate Contamination .... 49 2.4 Posting of MoSi2 ...... 49
III. EXPERIMENTAL PROCEDURE AND MATERIALS 52 3.1 Materials ...... 52 3.2 Ti-Si Diffusion Couple ...... 54 3.3 Pack Cementation Method .... 54 3.4 SOLGAS Calculations ...... 57
Vii 3.5 Oxidation Evaluation ...... 61 3.5.1 Isothermal Oxidation ...... 61 3.5.2 Cyclic oxidation ...... 62 3.6 Analysis, Characterization and Evaluation of Coatings ...... 66 3.7 Salt Coating Method of MoSi2 ...... 67 THE GROWTH AND DEVELOPMENT OF BORON- AND GERMANIUM-DOPED SILICIDE COATINGS: RESULTS AND DISCUSSION ...... 69 4.1 Titanium-Silicon Diffusion Couple ...... 69 4.1.1 Theory of Five-LayeredGrowth ...... 70 4.1.2 Multi-Layered Growth Global Growth Rate ...... 74 4.1.3 Growth Kinetics ...... 79 4.2 Growth of Undoped Silicide Coatings ...... 86 4.2.1 Thermodynamics of Siliconizing .... 87 4.2.2 Characterization of As-Coated S u b s t r a t e s ...... 92 4.2.3 Growth of Silicide Diffusion C o a t i n g s ...... 96 4.2.4 Global Growth Rates ...... 107 4.3 Growth of Titanium-Boride Diffusion Coatings .107 4.3.1 Thermodynamics of Boriding ...... ill 4.3.2 Kinetics of Boriding ...... Ill 4.4 Boron-Doped Silicide Coatings ...... 119 4.4.1 Development of B-Doped Silicide Coatings on CP-Titanium ...... 120 4.4.2 Growth Kinetics of B-Doped Silicide Coatings on CP-Titanium ...... 138 4.4.3 B-Doped Silicide Coatings on Ti-Al-Nb A l l o y s ...... 147 4.5 Germanium-Doped Silicide Coatings ...... 156 4.5.1 Ge-Doped Silicide Coatings on CP-Titanium ...... 157 4.5.2 Variation of Ge-Content...... 171 4.5.3 Growth Kinetics of Ge-Doped Silicide C o a t i n g s ...... 172 4.5.4 Ge-Doped Silicide Coatings on Ti-Al-Nb Alloys ...... 182 4.6 Ge-Doped and Undoped Molybdenum-Silicide C o a t i n g s ...... 193 4.6.1 Characterization of the Coatings . . . 194 4.6.2 Growth Kinetics of Ge-Doped Molybdenum Silicide ...... 199
viii LIST OF TABLES
TABLE PAGE 1. Compositions of pack powders (in wt.%) used in the development of a B-doped silicide coating. . . . 58 2. Compositions of pack powders (in wt.%) used to develop the Ge-doped silicide coating...... 59 3. Pack compositions used to produce boride coatings on CP-titanium (wt.%)...... 59 4. Compositions of pack powders (wt.%) that produced the best silicide coatings on CP-titanium and were used for further evaluation by isothermal and cyclic oxidation ...... 60 5. Pack composition (wt.%) used to produce Ge-doped silicide coatings where the ratio of Si to Ge was varied in the pack...... 60
6 . Summary of the different powder pack mixtures used to produce undoped and germanium-doped silicide coatings on CP molybdenum (wt.%) ...... 61 7. Standard Gibbs energy of formation (AG°[J/mole]) for the five compounds in the Ti-Si system [138] at 950, 1050 and 1150°C ...... 75
8 . Theoretical parabolic rate constants (kp [cm2/s]) for the growth of TiSi2 and Tisi in the multi-layered TiSi2/TiSi/Ti5Si4/TisSi;i/Ti3Si coating at 950, 1050 and 1150°C. Diffusion data from literature are usedin the calculation .... 75
9. Intrinsic parabolic rate constants (k„ [cm2/s]) measured by a titanium-silicon diffusion couple at 950, 1050 and 1150«C...... 81 10. Activation energies for the growth of the five silicide layers in a titanium-silicon diffusion c o u p l e ...... 82
ix 11. Standard Gibbs energies of formation normalized to one half mole of molecular fluorine for MgF2, A1F3, and CuF, at 1150°C [143], i.e.: (l/y)X + l/2F2 = (l/y)XFy ...... 88 12. The calculated partial pressures for equilibrium of silicon, MgF2 activator and A1203 filler at ...... 89
13. The sum of SiF3, SiF2 SiF and Si partial pressures for packs activated with MgF2, A1F3 and CuF2 calculated at 1150°C, 1050°C and 950°C ...... 89
14. Intrinsic parabolic rate constants (kp [cm2/s]) determined from the experimental data for a MgF2- activated pack at 950, 1050 and ll50dC ...... 97
15. Intrinsic parabolic rate constants (kp [cm2/s]) calculated from the experimental data of an A1F3- activated pack at 950, 1050 and 1150°C...... 98
16. Intrinsic parabolic rate constants (kp [cm2/s]) calculated from the experimental data for a CuF2- activated pack at 950, 1050 and 1150°C ...... 99 17. Activation energies (Q [J/mole]) for the growth of each layer of the multi-layered coating for MgF2-, A1F3- and CuF2-activated p a c k s ...... 100
18. Global growth rate (cm2/s) for the overall coating that are calculated from the intrinsic rate constant for MgF2-, AlF3- and CuF2-activated packs at 950°c, 1050°C and 1150°C...... 108
19. Global rate constants (cm2/s) for the overall growth of the titanium-silicon diffusion couple calculated from the intrinsic rate constants at 950°C, 1050®c and 1 1 5 0 ° C ...... 109
20. Global rate constants (cm2/s) for the overall growth rate of the five-layered silicide coatings measured at 950°c, 1050° and 1150°C for MgF2, A1F3 and CuF2 activators, and for the titanium-silicon diffusion couple ...... 110
21. The sum of BF2, BF, B and B2F4 partial pressures calculated for packs activated by MgF2 andA1F 3 . . 113
x 22. Intrinsic parabolic rate constants (kp [cm2/s]) determined for the growth of the dual-layer TiB2/TiB coating by packs activated with either MgF2 or A1F3 at 950°C, 1050#C and 1150<*C...... 115
23. Intrinsic parabolic rate constants (k. [cm2/s]) for TiB2 growth at 950°C, 1050°C, and 1150°C that are calculated from the grain boundary diffusion data reported by Samsonov and Latysheve [149] ...... 116 24. Activation energies (Q [J/mol]) determined from the intrinsic parabolic rate constants for the dual-layered growth of TiB2 and TiB by an A1F3- and MgF2-activated pack. Activation energy determined from the intrinsic rate constant for TiB2, which was calculated from the grain boundary diffusion data of Samsonov and Latysheve [149] . . .117 25. The activity of boron and Gibbs energy of formation (AG°[J/mol]) normalized to one mole of boron for SiB3, CrB2, TaB2 and TiB2 [ 1 5 0 , 1 5 1 ] ...... 129
26. The average thicknesses of the TiB2, TiSi2, TiSi, TisSi4, Ti5Si3 and Ti3Si layers that are produced by a pack comprised of either 7 wt.% Si, 6 % TiB2, 2% MgF2 and A1203, or 7 wt.% Si, 6 % TaB2, 2% MgF2 and A1203 at 1150°C for 12 h o u r s ...... 135
27. The average thicknesses ([pm]) of the TiB2 , TiSi2, TiSi, Ti5Si4, Ti5Si3 and Ti3Si layers grown by a pack comprised of either 7 wt.% Si, 6 % TiB,, 2% A1F3 and A1203, or 7 wt.% Si, 6% TaB2, 2% A1F3 and A1203 at 1150°C for 12 hours...... 136
28. The average thicknesses ([/m]) of the TiB2, TiSi2, TiSi, Ti5Si4, Ti5Si3 and Ti3Si layers grown by a pack comprised of either 7 wt.% Si, 6 % TiB2, 2% CuF2 and A1203, or 7 wt.% Si, 6% TaB2, 2% CuF2 and A1203 at 1150°C for 12 hours ...... 137
29. The average thicknesses ([/xm]) of the TiB2, Tisi2, TiSi, Ti5Si4, Ti5Si3 and Ti3Si layers grown by a pack comprised of either 7 wt.% Si, 6 % TiB2, 2% MgF2 and A1203, or 7 wt.% si, 6% TiB2, 2% CuF2 and A1203 at either 950°C for 6 hours, or 950°C for 28 hours ...... 140
xi 30. Intrinsic parabolic rate constants (kp[cm2/s)) measured for TiSi2, TiSi, TicSi4, Ti5Si3, and Ti3Si layers grown on a B-doped silicide coating by a pack comprised of 7 wt.% Si, 6 % TiB2, 2 % MgF2 and A120 3 at 950°C, 1050°C, and 1150°C...... 141
31. Intrinsic parabolic rate constants (kp[cm2/s]) determined for the TiB2 layer and the TiSi2, TiSi, T i 5Si4, Ti5Si3, and Ti3Si layers grown on a B-doped silicide coating by a pack comprised of 7 wt.% Si, 6 % TiB2, 2% A1F3 and A1203 at 950°C, 1050°C, and 1 1 5 0 ° C ...... 142
32. Activation energies (Q[J/mole]) determined from the intrinsic parabolic rate constants for the growth of B-doped silicide layers by packs composed of 7 wt%. Si, 6% TiB2 and A120 3 filler with either MgF2 or A1F3 a c t i v a t o r ...... 143
33. The average thickness ([nm]) for the TiB2, TiSi2 and TiSi layers for B-doped silicide coatings grown on Ti-22Al-27Nb (22-27) and Ti-20Al-22Nb (20-22) by 8 different pack treatments; (1) Si-TiB2/MgF2/Al203 at 1150«C for 12 hours, (2) Si-TiB2/MgF2/Al203at 950°C for 28 hours, (3) Si-TiB2/MgF2/Al203at 950°C for 6 hours, (4) Si-TaB2/MgF2/Al203 at 1150°C for 12 hours, (5) Si-TiB2/CuF2/Al203at 950°C for 28 hours, (6 ) Si-TiB2/CuF2/Al203at 950°C for 6 hours, (7) Si-TaB2/CuF2/Al203at 950°C for 28 hours, (8 ) Si-TaB2/CuF2/Al203at 950°C for.6 hours...... 157
34. The average thicknesses ([jxm]) for the Ti(Si,Ge)2, Ti(Si,Ge), Ti5 (Si,Ge)Ti5 (Si,Ge)3, and Ti3 (Si,Ge) layers grown on CP-titanium by a pack comprised of 12 wt.% Si, 6 % Ge, 2% MgF2, and Sic for three coating treatments: (1) 1150°c for 12 hours (2) 950°C for 28 hours (3) 950°C for 6 hours . . . .169 35. The average thicknesses {[/im]) for the Ti(Si,Ge)2, Ti(Si,Ge), Ti5 (Si,Ge)i, Ti5 (Si,Ge)3, and Ti 3 (Si,Ge) layers grown on CP-titanium by a pack comprised of 16 wt.% Si, 8% Ge, 2% A1F3, and Al203 for three coating treatments: (1) 1150°C for 12 hours (2) 950°c for 28 hours (3) 950°c for 12 hours . . . 169 36. The average thicknesses ([jm]) for the Ti(Si,Ge)2, Ti(Si,Ge), Ti5 (Si,Ge)4, Ti5 (Si,Ge)3, and Ti3 (Si,Ge) layers grown on CP-titanium by a pack comprised of 16 wt.% Si, 8 % Ge, 2% CuF2, and Al20 3 for three coating treatments: (1) 1150°C for 12 hours (2) 950°C for 28 hours (3) 950°C for 12 hours . . . 170
xii 37. The average thicknesses ([Mm]) for the Ti(Si,Ge)2, Ti(Si,Ge), Ti 5 (Si,Ge)4, Ti5 (Si,Ge)3, and Ti3 (Si,Ge) layers grown on CP-titanium by a pack comprised of Si, Ge, 2% MgF2, and Sic at 1150°C with three Si to Ge ratios: (1) 12 wt.% Si and 6 % Ge [2:1] (2) 12% Si and 12% Ge [1:1] (3) 6% Si and 12% Ge [1:2] ...... 173
38. The average thicknesses ([ Mm]) for the Ti(Si,Ge)2, Ti(Si,Ge), Ti5 (Si,Ge)4, Ti 5 (Si,Ge)3, and Ti3 (Si,Ge) layers grown on CP-titanium by a pack comprised of Si, Ge, 2 wt.% A1F3, and A120 3 at 1150°C for 12 hours with three Si to Ge ratios: (1) 16 wt.% Si and 8 % Ge [2:1] (2) 16% Si and 16% Ge [1:1] (3) 8 % Si and 16% Ge [1:2] ...... 174 39. The average thicknesses ([Mm]) for the Ti(Si,Ge)2, Ti(Si,Ge), Ti5 (Si,Ge)4, Ti5 (Si,Ge)3, and Ti3 (Si,Ge) layers grown on CP-titanium by a pack comprised of Si, Ge, 2 wt.% CuF2, and A1203 at 1150°c for 12 hours with three Si to Ge ratios: (1) 16 wt.% Si and 8% Ge [2:1] (2) 16% Si and 16% Ge [1:1] (3) 8 % Si and 16% Ge [1:2] ...... 175
40. Intrinsic parabolic rate constants (kp[cm2/s]) determined for the Ti(Si,Ge)2, Ti(Si,Ge), Ti 5 (Si,Ge)4, Tis(Si,Ge)3, and Ti3 (Si,Ge) layers grown by a pack comprised of 16 wt.% Si, 8 % Ge, 2% A1F3 and A1203 on CP-titanium at 1150°C, 1050°C and 950°C ...... 177
41. Intrinsic parabolic rate constants (kp[cm2/s]) determined for the Ti(Si,Ge)2, Ti(Si,Ge), Ti 5 (Si,Ge)4, Ti5 (Si,Ge)3, and Ti3 (Si,Ge) layers grown by a pack comprised of 12 wt.% Si, 6 % Ge, 2% MgF2 and Sic on CP-titanium at 1150°C, 1050°C and 950#C ...... 177
42. Activation energies (Q[J/mol]) determined for the growth of the Ti(Si,Ge)2, Ti(Si,Ge), Ti5 (Si,Ge)4, Tis(Si,Ge)3, and Ti3 (Si,Ge) layers by the MgF2- ana AlF3-activated packs ...... 177
43. The average thicknesses ( [Mm]) for the Ti(Si,Ge)2 and Ti(Si,Ge) layers grown on Ti-22Al-27Nb and Ti-22Al-27Nb by a pack comprised of 12 wt.% si, 6 % Ge, 2% MgF2, and sic at three coating treatments: (1) 1150°c for 12 hours on Ti-22Al-27Nb (2) 950°C for 28 hours on Ti-20Al-22Nb (3) 950°C for 6 hours on Ti-20Al-22Nb ...... 190
xiii 44. The average thicknesses ([jxm]) for the Ti(Si,Ge)2 and Ti(Si,Ge) layers grown on Ti-J22Al-27Nb and Ti-22Al-27Nb by a pack comprised of 16 wt.% si, 8% Ge, 2% AlF-j, and A120 3 at three coating treatments: (l) 1150®C for 12 hours on Ti-22Al-27Nb (1) 1150°C for 12 hours on Ti^20Al-22Nb (2) 950°C for 28 hours on Ti*^20Al-22Nb (3) 950°C for 6 hours on Ti-20Al-22Nb ...... 191
45. The average thicknesses {[/in]) for the Ti(Si,Ge)2 and Ti(Si,Ge) layers grown on Ti-22Al-27Nb and Ti-22Al-27Nb by a pack comprised of 16 wt.% Si, 8% Ge, 2% cuF2, and A120 3 at three coating treatments: (l) 1150°C for 12 hours on Ti-20Al-22Nb (2) 950°c for 28 hours on Ti-20Al-22Nb (3) 950°C for 6 hours on Ti-20Al-22Nb ...... 192 46. Elements present in the byproduct layers on the surface of MoSi2 coatings grown by NaF- and MgF?- activated packs (Pack# 1 and 3). Major: 15-60 at.%, Medium: 3-10 at.%, Minor 0.5- 2 at.%...... 198
47. The calculated equilibrium partial pressures for pack comprised of pure Si, pu re Ge, NaF activator and Sic filler at 1423 K ,200
48. Intrinsic parabolic rate constants (k_ [cm2/s]) determined from the experimental data for a Ge-doped silicide coating produced using a lOSi-lOGe/NaF/SiC pack at 950, 1050 and 1150°C. . . 206
49. Theoretical intrinsic parabolic rate constants (kp [cm2/s]) for the growth or a three-layer molybdenum-silicide coating at 950, 1050 and 1150°C that are calculated from diffusion data reported in the literature ...... 206
50. Intrinsic parabolic rate constants (kp [cm2/s]) for the growth of a three-layer Mo-Ge coating at 950, 1050 and 1150°C that are calculated from the reported rate constants for a Mo-MoGe2 diffusion couple [155], and the single-layer growth rate for MoGe2 reported in literature [155] 206
51. Activation energies (Q [J/mol] ) determined from the growth of each layer for the three-layered Ge-doped silicide coating, and calcula ted values for the rate constants for MoSi2, and Mo5Si3 and Mo3Si. . .208
xiv LIST OF FIGURES
FIGURE PAGE 1. Arrhenius plot of the parabolic rate constants for AI9O 3 , Cr20 3 and Si02 scales produced by the oxidation of pure materials oralloys [ 5 ] ...... 4 2. Phase diagram of the Ti-Si system [57]...... 14 3. Schematic illustration of the (a) circulation mechanism, and (b) condensation mechanism described by Levine and Caves [6 6 ] where circles are Al(l) and the triangles are AlF3 ( c ) ...... 23
4. Extension of the condensation mechanism (a) Mixed mechanism of A1 transport for a condensed activator with an activator and activator/source depleted zone, (b) three regions described for the condensed activator mechanism for the aluminizing of Ni by an A1F3 activator [85]...... 24 5. Log partial pressure as a function of temperature calculated for a pack containing a NaCl activator and pure Al, Cr and Si [35] ...... 29
6 . Phase diagram of the Mo-Si system [57]...... 33 7. Phase diagram for the Mo-Ge system[57] ...... 33
8 . Cyclic oxidation kinetics at 1370°C of undoped MoSi2 coatings compared with Ge-doped MoSi2 diffusion coatings on pure Mo [25]...... 34 9. Schematic of the (a) dual-layer oxide scale described by Yurek, Hirth and Rapp [110], and (b) the triple-layer oxide scale described by Hsu[111] . 45 10. Schematic of the experimental setup used to coat the workpieces by the pack cementation method. . . . 56 11. Schematic diagram the Cahn-171 thermogravimetric analyzer (TGA) used to monitor the weight change for isothermal oxidation ...... 63
xv 1 2 . Schematic illustration of the cyclic oxidation rigs used for this study ...... 65 13. Optical micrograph of the cross-section of the titanium-silicon diffusion couple following 14 hours at 1150°C. The five TiSi2/TiSi/TisSi4/Ti5Si3/Ti3Si layers are revealed by the Kroll etch...... 70 14. A schematic drawing of the five-layered silicide coating that is superimposed on the activity profile for silicon in the Ti-Si system at 950, 1050 and 1150°C [13B]...... 76 15. A linear regression fit to the plot of the thickness of the TiSi2, TiSi, Ti5Si4, T i 5Si3, and Ti3Si layers versus tne square root of time for the titanium-silicon diffusion couple 1150°C. The error bars are one standard deviation, and no error bar is shown for a standard deviation smaller than the plotting symbol...... 81 16. A linear regression fit to an Arrhenius plot of log intrinsic parabolic rate constant versUB inverse temperature for the TiSi2, TiSi, Ti5Si4 , Ti5Si3, and Ti3Si layers in the titanium-silicon diffusion couple. The error bars are the estimated standard deviation, with no error bar shown if the plot symbol is smaller than the standard deviation. . . . 82 17. A linear regression fit to an Arrhenius plot of the calculated diffusion coefficient versus inverse temperature for Tisi2, TiSi, Ti5Si4, Ti5Si3, and Ti3Si from the diffusion couple measurements .... 88 18. Thermodynamic calculation of the partial pressures of the silicon fluorides for pure silicon, MgP2 activator and A1203 filler at 1000K to 1600K .... 90 19. Thermodynamic calculation of the partial pressure of the silicon fluorides species of pack consisting of pure silicon, A1F3 activator and A1203 filler at 1000K to 1600K ...... 90 20. Thermodynamic calculation of the partial pressure of the silicon fluorides species of pack consisting of pure silicon, CuF2 activator and A1203 filler at 1000K to 1600K ...... 91
xvi 21. Quantitative EDS profiles of silicon in the cross- section of the multilayered silicide coating after 12 hours at 1150°C: (a) 10 wt% Si, 2 wt% MgF2 and A1,0, filler, (b) 10 wt% Si, 2 wt% A1F3 and A120 3 f i l l e r ...... 94
22. Optical micrograph of the cross-section of a silicide coating grown for 12 hours at 1150°C. The Kroll etch was used and the five TiSi 2/TiSi/Ti5Si4/Ti5Si3/Ti3Si layers are marked on the micrograph, (a) 10 wt% silicon, 2 wt% MgF* and A 1 20 3 pack, (b) 10 wt% silicon, 2 wt% A1F3 ana A 1 20 3 pack, (c) 10 wtt silicon, 2 wt% CuF2 and A 1 20 3 p a c k ...... 95
23. A linear regression fit to the plot of the thickness of the TiSi2, TiSi, TisSi4, Ti5Si3, and Ti3Si layers versus the square-root of time in a coating grown at 1150°C from a 10wt% silicon, 2 wt% MgF2 and A120 3 filler pack ...... 97
24. Plot of the thickness of the TiSi2, TiSi, Ti5Si4 , T i 5Si3, and Ti3Si layers versus the square-root of time for a coating grown at 1150°C with a pack composed of 10wt% silicon, 2 wt% A1F3 and A120 3 f i l l e r ...... 98
25. Plot of the thickness of the Tisi2, TiSi, Ti5Si4 , TisSi3, and Ti3Si layers versus the square-root of time for a coating grown at 1150°C with a pack composed of 10wt% silicon, 2 wt% CuF2 and A1203 f i l l e r ...... 99
26. A linear regression fit to an Arrhenius plot of log intrinsic parabolic rate constant versus inverse temperature for the TiSi2, TiSi, Ti 5Si4, T i 5Si3, and Ti3Si layers in a coating grown toy a MgF2- activated pack ...... 100
27. A linear regression fit of an Arrhenius plot of log intrinsic parabolic rate constant versus inverse temperature for the TiSi2, TiSi, Ti5si4, Ti5si3, and Ti3Si layers ofan AlF3-activated pack .... 101 28. A linear regression fit to the Arrhenius plot of log intrinsic parabolic rate constant versus inverse temperature for the TiSi2, TiSi, TicSi4, T i 5Si3, and Ti3Si layers in a coating formed by of a CuF2-activated p a c k ...... 102
xvii 29. Thermodynamic calculation of the equilibrium boron fluoride pressures for a pack comprising of B, MgF2 and A1203 at 1000 to 1GOOK...... 109 30. Thermodynamic calculation of the equilibrium boron fluoride pressures for a pack comprising of B, MgF2 and A1203 at 1000 to 1600K...... 110 31. Optical micrograph of the cross-section of a boride coating grown by a pack comprised of 4 wt.% B, 2% HgF2, and A1203 at 1150 °c for 12 hours...... 112 32. Optical micrograph for a boride coating grown by a pack comprised of 4 wt.% B, 2% A1F3, and A1203 at 1150°C for 12 hours. The TiB2 and TiB phases are marked in the cross-section...... 113 33. A linear regression fit for the plot of the average TiB2 and TiB thickness versus the square root of time for a MgFjj-activated pack at 950°C, 1050°C, and 1150°C. Tne error bars represent one standard d e v i a t i o n ...... 115
34. A linear regression fit for the plot of the TiB2 and TiB layer thickness versus the square root of time for an AlF3-activated pack at 950°C, 1050°C, and 1150°C. The error bars represent one standard d e v i a t i o n ...... 116 35. A linear regression fit for an Arrhenius plot of the intrinsic parabolic rate constants versus inverse temperature for the TiB2 and TiB layers formed by either a MgF2- or AlF3-activated p a c k ...... 117 36. Partial pressures of Si and B fluorides calculated for a pack comprised of pure Si, pure B, and a MgF2 activator at 1000-1600 K ...... 122
37. The Si and B activity in the Si-B system calculated at 1400K from the data of Dirkx and Spear [150] . . 122
38. Partial pressures for the Si and B fluorides calculated for a pack comprised of pure Si, a siB3 compound, and a MgF2 activator at 1000-1600 K . . . 123 39. The Si and B activity in the Ti-B systemcalculated at 1400K [151] ...... 125 40. Partial pressures for the Si and B fluorides calculated for a pack comprised of pure Si, a TiB2 compound, and a MgF2 activator at 1000-1600 K . . . 125
xviii 41. The polished cross-section of the B-doped silicide coating grown on CP-titanium toy a pack composed of 7 Wt.% Si, 6 % TiB2, 2% MgF2, and A1203 at 1150°C for 12 hours, (a) optical micrograpn, (to) optical micrograph of etched cross-section and (c) microprotoe profile of Si and B ...... 126 42. A SEM micrograph of the surface of a B-doped silicide coating grown on CP-titanium by a pack consisting of 7 wt.% si, 6 % TiB2, 2% MgF2 and A1203 at 1150°C for 12 hours...... 128 43. Partial pressures of the Si and B fluorides calculated for a pack composed of pure Si, aCrB 2 compound, and a MgF2 activator at 1000-1600 K .. . 128
44. Partial pressures for the Si and B fluorides calculated for a pack comprised of pure Si, a TaB2 compound, and a MgF2 activator at 1000-1600 K . . . 129 45. Optical micrograph of the polished cross-section of a B-doped silicide coating on CP-titanium grown by a pack comprised of 7 wt.% Si, 6 % crB2, 2% MgF2 and Al2o3 at H50°c for 12 hours...... 130
46. Optical micrograph of the polished cross-section of a B-doped silicide coating on CP-titanium grown by a pack comprised of 7 wt.% Si, 6 % TaB2, 2% MgF 2 and A1203 at 1150°C for 12 hours...... 130 47. The equilibrium vapor pressures for the Si and B fluorides calculated for a pack comprised of pure Si, a TiB2 compound, and an A1F3 activator at 1000-1600 K ...... 134
48. Optical micrograph of the polished and etched cross-section for a B-doped silicide coating formed by a pack comprised of 7 wt.% Si, 6 % TiB2,2% A1F3 and A120 3 at 1150°C for 12 hours ...... 134
49. The equilibrium partial pressures of the Si and B fluorides calculated for a pack comprised of pure Si, a TaB2 compound, and an A1F3 activator at 1000-1600 K ...... 135
50. The equilibrium partial pressures of the Si and B fluorides calculated for a pack comprised of pure Si, a TiB2 compound, and a CuF2 activator at 1000-1600 K ...... 136
xix 51. The equilibrium partial pressures of the Si and B fluorides calculated for a pack comprised of pure Si, a TaB9 compound, and a CuF, activator at 1000-1600 K ...... 137 52. Linear regression fit to a plot of thickness versus time for TiSi2, TiSi, TisSi4, Ti5Si3, and Ti3Si layers formed on a B-doped silicide coating by a pack comprised of 7 wt.% Si, 6 % TiB2, 2% MgF2 and A1203 at 1150°C. The error bars are one standard deviation...... 141 53. Linear regression fit to a plot of thickness versus time for TiSi2, TiSi, Ti5Si4, T i 5Si3, and Ti3Si layers formed on a B-doped silicide coating by a pack comprised of 7 wt.% Si, 6 % TiB2, 2% A1F3 and A1203 at 1150°C. The error bars are one standard deviation...... 142 54. Linear regression fit to an Arrhenius plot of the intrinsic parabolic rates for the TiSi2, TiSi, Ti5Si4, TieSi3, and Ti3Si layers formed by a HgF2- activated pack ...... 143 55. Linear regression fit to an Arrhenius plot of the intrinsic parabolic rates for the TiSi2, TiSi, Ti5Si4, Ti 5Si3, and Ti3Si layers formed by an AlF3-activated pack ...... 144
56. Linear regression plots for the TiB2 layer on the B-doped silicide coating grown by a pack comprised of 7 wt.% Si, 6 % TiB2, 2% A1F3 and Al203; (a) thickness versus the square root of time, and (b) Arrhenius plot for the intrinsic rate constants...... 145 57. B-doped silicide coating grown on Ti-22Al-27Nb by a pack comprised of 7 wt.% Si, 6 % TiB2, 2% MgF2 and A120. at 1150°C for 12 hours, (a) Optical micrograph of tne polished and etched croBs-section and (b) EDS profile of Si, Nb and A1 in the coating . . 151
58. B-doped silicide coating grown on Ti-22Al-27Nb by a pack comprised of 7 wt.% Si, 6 % TiB2, 2% MgF2 and A120 3 at 1150°C for 12 hours, (a) SEM/BSE micrograph of a polished cross-section and (b) microprobe profile of B, Si, Nb and A1 in the coating ...... 151
XX 59. Optical micrograph of the corner of a B-doped silicide coating grown on a Ti-22Al-27Nb alloy by a pack comprised of 7 wt.% Si, 6 % TiB2, 2% MgF, and A120, at 1150®C for 12 hours, which shows-the convolution ...... 152 60. B-doped silicide coating grown on Ti-20Al-22Nb by a pack comprised of 7 wt.% si, 6% TiB2, 2% MgF2 and A1,03 at 950®C for 6 hours, (a) SEM micrograph of a polished and etched cross-section and (b) EDS profile of B, Si, Nb and A1 in the coating ...... 154 61. Optical micrograph of the polished cross-section of a B-doped silicide coating grown on a Ti-48Al-2Mn-2Nb by a pack comprised of 7 wt%. Si, 6 % TiB2, 2% MgF2, and A1203 at 1150°C for 12 h o u r s ...... 158
62. Thermodynamic calculation of the Si and Ge fluorides for a pack composed of Si, Ge, MgF2 and Sic at 1000-1600 K ...... 158 63. Ge-doped silicide coating formed by a pack comprised of 12 wt.% Si, 6 % Ge, 2% MgF2 and Sic at 1150°C for 12 hours, (a) Optical micrograph and (b) EDS profile for Si and G e ...... 159 64. Thermodynamic calculation of the partial pressures of the Si and Ge fluorides for a pack comprised of (a) Si, Ge, and A1F3 and (b) Si, Ge, and A1F3 at 1000-1600 K ...... 164 65. Ge-doped silicide coating (a) Optical micrograph for the polished cross-section of a formed by a pack comprised of 16 wt.% Si, 8 % Ge, 2% A1F3 and A120 3 at 1150°C for 12 hours, and (b) SEM micrograph of a Ge-doped silicide coating grown by a pack comprised of 16 % Si, 8 % Ge, 2% CuF2 and A120 3 at 1150°C for 12 hours ...... 165
6 6 . A convolution in a Ge-doped silicide coating grown at the corner of a CP-titanium workpiece by a pack composed of 16 wt.% Si, 8 % Ge, 2% CuF2 and A120 3 at 1150°C for 12 hours ...... 166 67. EDS profile of Si, Ge and A1 for the Ge-doped silicide coating (a) formed by a pack composed of 16 wt.% Si, 8 % Ge, 2% AlF-j and A1 203 at 1150°C for 12 hours, and (b) formed by a pack composed of 16 wt.% Si, 8 % Ge, 2% CuF2 and A120 3 at 1150°C for 12 h o u r s ...... 168
xxi 68. EDS profile of Ge for the Ge-doped silicide coating formed by a pack comprised of Si, Ge, 2 wt.% MgF2 and Sic at 1150°C for 12 hours with three ratios of Si to Ge: (1) 12% Si and 6 % Ge [2:1], (2) 12% Si and 12% Ge [1:1], and (3) 6% Si and 12% Ge [1:2] . .173 69. EDS profile of Ge for the Ge-doped silicide coating formed by a pack comprised of Si, Ge, 2 wt.% A1F3 and A120 3 at 1150*C for 12 hours with three ratios of Si to Ge: (1) 16% Si and 8 % Ge [2:1], (2) 16% Si and 16% Ge [1:1], and (3) 8% Si and 16% Ge [1:2] . .174 70. EDS profile of Ge for the Ge-doped silicide coating formed by a pack comprised of Si, Ge, 2 wt.% CuF2 and A1203 at 1150°C for 12 hours with three ratios of si to Ge: (l) 16% Si and 8 % Ge [2:1], (2) 16% Si and 16% Ge [1:1], and (3) 8 % Si and 16% Ge [1:2] ...... 175
71. Linear regression fit to a plot of thickness versus the square root of time for Ge-doped silicide coatings grown at 1150°C by packs comprised of (a) 12 wt.% Si, 6% Ge, 2% MgF2, and Sic, and (b) 16% Si, 8% Ge, 2% AIF3 and A1203 ...... 176
72. A linear regression fit to an Arrhenius plot of log intrinsic rate constant versus inverse temperature for the Ti(Si,Ge)2, Ti(Si,Ge), Ti5 (Si,Ge)4, Tie(Si,Ge>3 , and Ti3 (Si,Ge) layers grown by (a) MgF2-activated pack and (b) AlF3-activated p a c k ...... 178 73. SEM micrograph of the polished cross-section of a Ge-doped silicide coating grown on a Ti-22Al-27Nb alloy by a pack comprised of 12 wt.% si, 6 % Ge, 2 % MgF2 and sic at ll50°c for 12 hours...... 183
74. EDS profiles for Ge-doped silicide coating grown on Ti-22Al-27Nb by a MgF2-activated pack at 1150°C for 12 hours, (a) Si, Ge, A1 and Nb profile and (b) Si, Al, and Nb profile ...... 184
75. Optical micrograph of the polished cross-section of a Ge-doped silicide coating grown on a Ti-22Al-27Nb alloy by a pack comprised of 16 wt.% Si, 8 % Ge, 2% AIF3 and Al203 at 1150°C for 12 hours ...... 183
xxii 76. EDS profiles of Ge-doped silicide coating grown on Ti-22Al-27Nb by an AlF3-activated pack at 1150°C for 12 hours, (a) Si, Ge, A1 and Hb profile and (b) Si, Al, and Mb profile ...... 186 77. SEM micrograph of the polished/etched cross-section of a Ge-doped silicide coating grown on a Ti-20Al-22Nb alloy at 950°C for 6 hours by a pack comprised of (a) 16 wt.% Si, 8 % Ge, 2% A1F3 and A1 203, and (b) 16 wt.% Si, 8 % Ge, 2% CuF2 and A1 20 3 ...... 189
78. EDS Si, Ge, Al and Nb profiles for a Ge-doped silicide coating grown on Ti-20Al-22Nb by a MgF2-activated pack at 950°C for 6 ho u r s ...... 190 79. EDS Si, Ge, Al and Nb profiles for a Ge-doped silicide coating grown on Ti-20Al-22Nb by an AlF3-activated pack at 950°C for 6ho u r s ...... 191 80. EDS Si, Ge, Al and Nb profile for a Ge-doped silicide coating grown on Ti-20Al-22Nb by a CuF2- activated pack at 950°C for 6 h o u r s ...... 192
81. cross-sections of a Ge-doped MoSi2 coating grown by a pack composed of lOSi-lOGe, 2NaF and Sic (weight%) at 1150°C for 12 hours, (a) SEM micrograph showing Mo(Si,Ge)2 and MOc(Si,Ge)3, and (b) SEM micrograph showing the thin layers of Mos(Si,Ge)3 and Mo3 ( S i , G e ) ...... 196
82. EDS profile of germanium in the three-layered Mo(Si,Ge)2/Mo5 (Si,Ge)3/Mo3 {Si,Ge) coating grown by a NaF-activated pack at 1150°C for 12 hours .... 197 83. SEM micrograph of the surface of a Ge-doped silicide coating grown by a NaF activated pack. The protrusions marked nan are sodium r i c h ...... 197 84. High-magnification SEM micrograph of the byproduct salt-oxide layers at the surface of MoSi2, (a) coating grown by a NaF-activated pack and (b) a MgF2-activated p a c k ...... 198
85. Linear regression fit to a plot of layer thickness versus the square-root of time for Ge-doped coatings grown by NaF-activated pack at 1150, 1050 and 950°C for (a) Mo(Si,Ge)2, (b) Mo5 (Si,Ge}3 , and (c) Mo3 (Si,Ge)...... 205
xxiii . Arrhenius plot of the intrinsic parabolic rate constant for (a) the experimentally measured Mo(Si,Ge)2 outer layers compared with the calculated intrinsic rate constants for MoSi2 [94,118,119}, and (b) measured intrinsic rate constants for the inner Mos(Si,Ge)3 and Mo 3 (Si,Ge) layers compared with calculated rates for the Mo 5Si3, M o 3Si, Mo 5Ge3 and Mo3Ge layers [115,116,155]...... 207
xxiv ABSTRACT
A procedure for codepositing either boron and silicon, or else germanium and silicon in a single reaction/processing step by a fluoride-activated, pack cementation method has been developed. The boron- or germanium-doped titanium-silicide diffusion coatings were formed on commercially pure (CP) titanium, Ti-22Al-27Nb, and
Ti-20Al-22Nb. The growth rates of undoped, B- and Ge-doped multi layer coatings were controlled by solid-state diffusion for less stable activators, and gas-phase diffusion for more stable activators. The boron was concentrated at the surface of the coating as a TiB2 layer, while the Ge was distributed within the coating.
The isothermal oxidation kinetics of the boron- and germanium-doped silicide coatings on CP-titanium and Ti-
22Al**27Nb are slower than pure TiSi2, and slightly faster than silicon oxidation. The activation energies for oxidation at 700-1000°C are equivalent to silicon, but a different mechanism was observed at 700-500°C due to the formation of mixed Ti02-Si02 scale.
xxv Slow cyclic oxidation kinetics were observed for the B- and Ge-doped silicide coatings on CP-titanium at 800-500°C. No substrate contamination was detected by microhardness, which demonstrates the coatings were protective. The coatings did not protect CP-titanium at 1000-900°C because of thermal stresses and the allotropic phase transformation
of titanium. Slow cyclic oxidation kinetics were observed for the B- and Ge-doped silicide coatings grown on Ti-22Al-27Nb and Ti-
20Al-22Nb at 1000-500°C and no substrate contamination was detected by microhardness. The growth of thick coatings resulted in convolutions at the corners, and the coatings spalled during low-temperature cyclic oxidation. The thin
coatings were not protective in cyclic oxidation at 1000°C. However, a medium thickness B- and Ge-doped silicide coating was protective at 1000-500°C.
The byproduct salt layer from the growth of MoSi2 diffusion coatings by a NaF-activated pack cementation method prevented accelerated, low temperature oxidation (Posting). The application of NaF layers stopped the pasting of bulk MoSi2, which are otherwise susceptible to posting. The rapid formation of a Na-silicate layer passivates HoSi2 and prevents the posting reaction.
xxvi CHAPTER I INTRODUCTION
1.1 REQUIREMENTS OF AEROSPACE MATERIALS The performance and durability of aerospace systems is generally limited by the materials of construction. The most stringent material specifications for an aerospace vehicle are in the propulsion system. A material must be lightweight, possess good high-temperature strength and creep resistance, impact and damage tolerance, fatigue resistance, ductility and environmental durability. None of the currently used engine materials meet all of the above criteria. Ni-base superalloys are the most widely used material systems for aircraft engines, and have been the workhorse alloys of the aerospace industry for over twenty years [1,2]. These alloys are ductile and have excellent high-temperature strength, creep, fatigue and impact resistance. The only limitations of Ni-base superalloys are the high density and poor oxidation resistance at high temperature (>700°C). The oxidation resistance of superalloys is improved by the application of a protective aluminide diffusion coating, which has proven to be a viable solution for over twenty years [3- 5]. However, the temperature requirements for advanced
1 propulsion technologies exceed the upper limit for Ni-base superalloys (-1000°C). Furthermore, the high density of superalloys creates efficiency and weight problems that limit the development of more powerful engines, which are required for large civilian aircraft. Many researchers are currently working on this problem by developing alternative material systems to replace the Ni-base superalloys [2].
1.2 TXTANXUM-BA8E ALLOYS One important criterion for the selection of a new aerospace material is the specific strength, which is the yield stress or elastic modulus divided by the density. The specific strength is a ration of the strength to weight provided by a material system. The density of Ti-base alloys is one half that of the superalloys, and the specific strength for many Ti alloys exceeds that of the superalloys [6]. Titanium-alloys possess good ductility and high-temperature strength up to 600°C, which is below the temperature limit for superalloys. However, alloys that are based upon a2-Ti3Al, orthorhombic-Ti2AlNb, and y~TiAl intermetallic compounds can be used at higher temperatures, and the specific strengths rival or exceed those of conventional superalloys [2,7,8].
Since the density of Ti-base alloys 1b much lower than superalloys and improvements in high-temperature strength have been achieved, the amount of Ti used in the production of aircraft has steadily increased over the last twenty years The most important limitation to the use of Ti-base alloys at high temperature is the environmental durability. A non-protective Tio2 scale forms on Ti during high- temperature oxidation, which has a rapid rate for oxygen diffusion and does not prevent the inward penetration of oxygen [10]. The base metal is consumed at a fast rate by the diffusive conversion to an oxide scale. Titanium has a large solubility for oxygen, and the base metal beneath the oxide can absorb up to 30 at.% oxygen [10]. The dissolution of oxygen by Ti results in severe embrittlement, which degrades the mechanical properties [11]. Alloys based upon Ti-rich
interroetallic compounds such as a2-Ti3Al, orthorhombic- Ti2AlNb, and y-TiAl also undergo a rapid rate of high- temperature oxidation and embrittlement by oxygen dissolution [7,8,12-16]. Since these intermetallic-base alloys possess a limited amount of room-temperature ductility, embrittlement by oxygen dissolution is a more serious issue. Two approaches for preventing the rapid rate of oxidation and embrittlement of Ti-base alloys are alloying or application of a protective coating. The alloy or coating must form a continuous A1203, Cr2o3 or Si02 scale during oxidation at high temperature to provide effective long-term protection. Figure 1 is an Arrhenius plot of the parabolic
rate constants for A1203, Cr203 and Si02, which shows the oxidation kinetics for alloys or compounds that form these scales are very slow [5], However, Cr203 forming alloys and 4
T.CC) 1200 11001000 900 000
CrjOj polycryttollin* rong* rfporltd by Lllltrud t Koftlod
AIjO, .polycryttollint ,.u rong* r*port*d by M. Hindom 1 Whlltl* SI 0] amorphous tattle* Deal i W illiams
,*>«! lotllc*
6 7 0 9 10 11 12 10 Vt,h-'
Figure 1: Arrhenius plot of the parabolic rate constants for Cr2°3 am* siC>2 scales produced by the oxidation of pure materials or alloys [5]. 5
coatings are not used at temperatures above 900#C because chromia is degraded by evaporative oxidation [17], Cr203 + (3/2)02 *• (2)Cr03 (1.1) Therefore, the ideal alloy or coating for Ti-base alloys should result in the growth of a protective A1203 or Si02 scale. Binary Ti-Al and Ti-Si alloys require either an Al content exceeding 50 at.% or a Si content exceeding 37.5 at.% for an A1203 or Si02 scale to be thermodynamically stable at 700-1300°C [18,19]. Aluminum or silicon contents that exceed these optimistic thermodynamic predictions are required to the support the growth kinetics of A1203 or Si02 scales [18-21]. Most binary Ti-Al or Ti-Si alloys that contain sufficient alloying additions to form an A1203 or Si02 scale cannot be used as a structural materials because of inadequate strength at high temperature or low ductility at room temperature. Large amounts of ternary or quaternary alloying additions are required to increase the chemical potential for Al or Si such that A1203 or Si02 are stable, and these alloys generally have poor mechanical properties. Except for some limited success with y-TiAl alloys [22], no Ti alloy has an acceptable amount of room temperature ductility and oxidation resistance at high temperature. Therefore, the most viable method for protecting Ti alloys against oxidation at high temperature is a coating treatment that enriches the surface of the base metal with Al or Si. 6
1.3 PROTECTION OF NIOBIUM ALLOYS The high-temperature oxidation of Nb-base alloys results in the same problems observed for Ti; Nb-base alloys form a non-protective Nb2Os scale, undergo a rapid rate of oxidation at high-temperature, and are embrittled by the dissolution of oxygen [10]. Niobium can be alloyed to improve the high- temperature oxidation resistance, but the mechanical properties of these alloys are not desirable [23,24]. Thus a protective coating is required for the practical application of Nb-base alloys at high temperature. Since the application temperatures for Nb alloys exceed the maximum temperature of superalloys, there is considerable interest in the development of a protective coating [25-28]. Recent work has demonstrated that a two-step coating process can provide excellent cyclic oxidation resistance to Nb alloys: (l) preliminary deposition of a Mo or Mo-W layer by physical vapor deposition (PVD), (2) partial diffusive conversion of the Mo or Mo-W layer into a
Ge-doped MoSi2 coating by a halide-activated pack cementation (HAPC) method [25-27]. One important limitation to the practical application of MoSi2 is the rapid, destructive form of low-temperature oxidation known as "pasting" [29-31]. The excellent high- temperature oxidation resistance of MoSi2 results from the exclusive formation of a protective Si02 scale [31-33]. At low-temperature (600-300°C) a dual-layer scale of Si02 and Mo- oxides form [29-31]. The rapid, voluminous formation of the 7
Mo-oxides results In accelerated crack growth at pre-existing cracks, pores and grain boundaries, which leads to the disintegration of MoSi2 (posting). Although the mechanisms and theory of pesting have been studied by many researchers, no practical solution has been discovered for this problem. Mueller et al. [26] did not observe pesting for
Mo(Si,Ge)2 and (Mo,W)(Si,Ge)2 diffusion coatings that were grown by the HAPC method after 192 hours of isothermal oxidation at either 500°C, 600°C or 700°C. However, an excess of 500 hours may be required to initiate the pesting reaction [31]. One objective of this work is to study the pesting resistance of MoSi2 and Mo(Si,Ge)2 diffusion coatings grown by the HAPC method.
1.4 PACK CEMENTATION COATINGS Hot dipping, electroplating, plasma spraying, sputtering, slurry fusion, thermal evaporation, ion implantation, chemical vapor deposition and pack cementation are some of the methods used to protect aerospace materials. These coating techniques are basically two separate processes: (l) physical vapor deposition (PVO) techniques such as sputtering and plasma spray, or (2) chemical vapor deposition (CVD) techniques such as pack cementation. Since large vacuum chambers and expensive equipment are required to produce the PVD coatings, these coatings are expensive. A post-coating heat treatment is usually required to achieve coating/substrate interdiffusion that improves the adhesion of these coatings. PVD methods are "line of sight” processes, and coating complex •# shapes is either difficult or impossible. CVD coatings are produced at high temperature by the reaction of gas-phase vapors to achieve deposition at the surface of the workpiece. Interdiffusion between the coating and substrate during the coating treatment results in excellent adherence for these coatings. A uniform coating will be deposited if the gas- phase vapors can reach all regions of the workpiece. However, the production of coatings with two or more elements by CVD methods is more difficult than by PVD. Since advanced coatings reguire multicomponent compositions, PVD is becoming a more popular method for coating aerospace materials, despite the increased cost. To deposit two or more elements in a single step by a CVD method, the thermodynamics of the system must be considered [34,35]. Processing variables such as temperature, total pressure, vapor pressures of reactive species, activity of source powders, carrier gas, deposition environment and substrate must be precisely controlled to obtain the correct "processing window", which may be very narrow in some cases [34,36]. Since coatings produced by CVD methods such as pack cementation are cheaper, more adherent, and easier to produce than PVD, a real opportunity exists for the commercialization of these processes. The halide-activated pack cementation (HAPC) method was first demonstrated by Van Aller [3] for the growth of aluminide coatings on Fe-base alloys. The process has been extensively commercialized, and is one of the most simple and economical means for applying a protective coating. The HAPC method involves embedding the workpiece to be coated into pre mixed powders that are sealed inside a refractory container called a retort. The three components of a powder pack are: (1) a masteralloy comprised of elements that are to be deposited at the surface of the workpiece, (2) a halide salt activator and (3) a relatively inert filler. The retorts are loaded into a furnace and shielded from oxidation during the isothermal heat treatment by an inert or reducing gas. The HAPC method is considered to be an in-situ CVD technique, because volatile halide species of the masteralloy are generated during the isothermal anneal by the decomposition of the halide activator and reaction with the masteralloy powder: Me(s) + z/yAXy(l,g) — > MXe(g) + A(l,g) (1.2) where Me is a masteralloy element, A is a cation in the halide salt and X is the halide anion. The difference in thermodynamic activity between the masteralloy powder and the workpiece drives gas-phase diffusion of the metal halide vapors to the surface of the workpiece. These metal halide vapors serve to transport the masteralloy element(s) to the surface to ach'ieve deposition and interdiffusion with the workpiece, while the depleted halide vapors recirculate back into the pack to react again with the masteralloy and repeat the cycle. The deposition of the masteralloy element(s) at the surface of a Ti workpiece occurs by exchange (1.3), 10
dissociation (1.4), disproportionation (1.5), or an alkali vapor species (1.6) reaction, MeXy(g) + Ti(s) * TiXy(g) + Me(s) (1.3) MeXy(g) = Me(s) + (y/2)X2(g) (1.4) (y+l)MeXy (g) = Me(s) + (y)MeFy+1(g) (1.5) A(v) + MeFy(g) + Me(s) + AXy(l,g) (1.6) The diffusion of the masteralloy element(s) into the workpiece forms an alloy if some mutual solubility exists, or otherwise a compound grows at the surface if the solubility limit is exceeded. In either case, the surface of the workpiece is converted to the coating. Additionally, solid-state
diffusion/reaction of the masteralloy element(s) with the workpiece maintains an activity gradient within the pack to continually drive gas-phase diffusion to the surface. The primary goal of this work is the development and evaluation of boron- and germanium-doped silicide coatingB for Ti alloys using a HAPC method. The purpose of the B or Ge addition is to dissolve into and improve the Si02 scale that forms on the silicide coating during oxidation at high temperature. Boron and germanium are both glass formers that increase the fluidity of Si02 to heal cracks at lower temperature [37,38). The higher coefficients of thermal expansion (CTE) for a Si02-B203 or Si02-Ge02 glass compared to pure Si02 minimizes spalling of the protective oxide scale upon thermal cycling, which results from large CTE differences between Si02 and the workpiece [39,40]. Since the oxidation kinetics for B and Ge are much faster than Si02, these elements serve as a fast growing transient oxide that protects the base coating, seals cracks, and are dissolved into Si02 [41,42]. Since the desired coating for titanium must prevent the penetration of oxygen and other contaminants into the titanium alloy substrate, the improvements to the oxide scale provided by the boron or germanium addition are required for optimum coating performance. The primary objectives of this thesis are: 1. Develop and optimize the HAPC method for producing B- or Ge-doped silicide coatings on Ti and Ti-aluminide alloys. 2. Determine the mechanisms and kinetics for the growth of undoped, and B- and Ge-doped silicide coatings on CP-
titanium. 3. Evaluate the performance of the coatings by isothermal
and cyclic oxidation. Determine if the coatings were effective barriers against substrate contamination. 4. Characterize the mechanism of high-temperature oxidation
for the B- and Ge-doped silicide coatings on Ti and Ti- aluminide alloys. CHAPTER II LITERATURE REVIEW
2.1 Halide-Aotivated Pack Cementation
2.1.1 Pack Siliconizing Compared to pack aluminizing and chromizing, little information is known about the growth or performance of Si- rich coatings. Therehas been considerable effort to produce
Si-rich coatings on Fe-base alloys, but a non-protective, porous coating was always produced [43,44]. Since the purpose of siliconizing iron is to produce aqueous corrosion resistance, a porous coating cannot be used. Silicide coatings have been applied to Ni-base alloys by pack cementation [45]. However, there are several low-temperature eutectics in the nickel-silicon system [46], which' excludes the possible use of a silicide coating on Ni-base alloys at high temperature. Additionally, these silicide compounds are very brittle and thermal cycling resulted in failure of these coatings [45]. Most metals and alloys have little solubility for Si, and the deposition of Si by a HAPC method results in the growth of a silicide compound. The most important limitation to the
12 13 practical use of silicide coatings is the lack of ductility or resistance to cracks that result from the coefficient of thermal expansion (CTE) mismatch [47,48]. However, Mueller et al. [25,26] demonstrated that Ge-doped MoSi2 silicide coatings provide excellent cyclic oxidation resistance to Nb at 1370°C and 1540°C. Additional work at United Technologies showed that plasma sprayed MoSi2-base coatings afforded excellent cyclic oxidation resistance to Nb alloys [28]. One reason for the excellent cyclic oxidation resistance is that the CTE for MoSi2 (a»8.5X10"6) is approximately the same as that for Nb (a«7.8X10“6) [27]. The siliciding of Ti by a HAPC method was invented by Yatema and Kluz in 1962 [49]. Further studies by other authors have shown that silicide coatings grown on Ti by the HAPC method provide good oxidation resistance at high temperature [50-55]. The coatings generally consist of multiple layers of Ti-silicide compounds. The outer layer is the Si-rich TiSi2 compound, which forms a slow growing Si02 scale during high-temperature oxidation [56]. The Ti-Si phase diagram in Fig. 2 shows five intermetallic compounds in the Ti-Si system, which are TiSi2, TiSi, Ti5Si4, Ti5Si3 and Ti3Si [57]. Thermodynamic calculations predict that a slow growing Si02 scale is stable on every silicide phase except Ti3Si [18,19]. Abba, Galerie and Caillet [58] used Ti5Si3 coatings to protect Ti and observed slow oxidation kinetics with a mixed Si02/Ti02 scale. 14 Weight Percent Silicon
m u '
iooo (si)—
*0 M IN Atomic Percent SiliconTI SI
Figure 2. Phase diagram of the Ti-Si system [57].
The formation of a continuous Ti02 scale would result in fast oxidation kinetics, and the results of Abba et al. [58] agree with the thermodynamic stability predictions for Si02. Since almost all the silicide phases stabilize a slow growing Si02 scale, limited interdiffusion between the Ti substrate should not degrade the effectiveness of the silicide coating. Furthermore, crack penetration through the outer silicide layers can be deflected by the lower silicide layers, which also form a protective Si02 scale. Unlike A1203, Si02 is an amorphous scale that can flow and heal cracks at high temperature. Therefore, silicide coatings are a viable approach for protecting Ti against oxidation scaling and 15 matrix contamination at high temperature.
2.1.2 Pack Boriding Boride compounds have a high hardness and excellent wear resistance, and there has been some interest in forming boride coatings by the HAPC method. Krzyminski [59] showed that the boride coatings formed on refractory metals possess excellent wear properties, singhal [60] borided various steels by the HAPC method, and demonstrated that the hardness of the FeB
coating exceeded that of carburized, nitrided or chrome plated steels. However, the low ductility of FeB was thought to limit wide application of the coating. In a separate study, singhal [61] borided Ti by a HAPC method, and showed that the erosion resistance was superior to any coated steel. The coating consisted of an outer TiB2 layer, which had a uniform thickness, and an inner layer of TiB with a non-uniform thickness. The growth kinetics for the dual-layered boride coating are extremely slow, with a 25 /mi thick coating produced by a 4B hour heat treatment at 1010°C on pure Ti. The growth kinetics for the boride coating on a Ti-6A1-4V alloy were even slower, with a 20 /m coating observed after the same coating treatment. 2.1.3 Pack Chromizing and Aluminising The chromizing process was invented by Kelly in 1923 [62] to protect plain carbon and low alloy steels from aqueous and high-temperature corrosion. As previously mentioned, Van Aller invented aluminizing in 1911 [3], The majority of 16 coatings produced by the HAPC method are aluminized or chromized, and the use of these coatings are well documented. Most of the studies on aluminizing have focused on coating Ni-base superalloys. The aluminizing process is classified as either a "high activity" or "low activity" coating, which is based on the coating microstructure and the Al source in the pack [63-65]. High activity aluminide coatings are formed by using a pure Al source at 700-900°C, which produces a high Al gas-phase flux. The solid-state diffusional growth of "high activity" coatings is dominated by inward transport of Al to form a Ni2Al3 layer. The Ni2Al3 coating contains a dispersion of cr- and Mo-rich precipitates, and dissolved alloying elements from the superalloy workpiece. A homogenization anneal is used to convert the Ni2Al3 phase into an Al-rich NiAl phase by the inward diffusion of Al. Three distinct zones are observed after the diffusion anneal: (1) an outer Al-rich NiAl phase with precipitates of Cr, Mo and substrate phases, (2) a NiAl zone with Cr, Mo, Ti and Co in solution, and (3) a NiAl matrix with refractory metal carbide and sigma phase precipitates. The second type of aluminizing treatment is a "low activity" process, which utilizes an Al-base masteralloy powder with a lower Al activity in the powder pack. Since the Al halide vapor pressures are reduced by using a masteralloy, a higher coating temperature (1000-1150°C) is required to achieve significant coating growth rates. The lower chemical 17 ' potential for Al in the gas phase results in the formation of a Ni-rich NiAl phase. A complex interdiffusion zone is formed between the baBe alloy and NiAl coating by the outward diffusion of Ni, and it is composed of block-shaped precipitates of refractory metal carbides (Ti, Nb, Mo, W, Hf, Ta) and rod-shaped sigma phase particles that contain Cr, Mo and Co. The outward diffusion of Ni results in the entrapment of pack particles in the coating. Since the pack particles degrade the performance of the coatings, in practice, the workpieces are physically isolated from the pack by an "above pack" process to eliminate pack entrapment. The chromizing of steels results in two different coatings, which are based on the amount of carbon in the steel workpiece [67,68]. The microstructure of coatings produced on low carbon steels (< 0.2 wt.%) consists of a a-(Fe-Cr) solid solution layer with carbide precipitates along the grain boundaries. A continuous surface layer of Cr-carbide results from chromizing steels with high carbon contents (> 0.2 wt.%). Chromizing is also used to improve the hot corrosion resistance of Ni-base superalloys. A Cr-carbide layer forms on carbon containing Ni- and Co-base superalloys. These alloys are coated at 1050°C or higher to permit carbide dissolution and increase the amount of Cr in the workpiece. 2.1.4 codeposition in a single Processing step
The oxidation resistance provided by a ternary alloy is generally superior to a binary alloy [35,71]. For example, Kanthal™ electric furnace heating elements are a Fe-20Cr-
4.5A1 alloy that derives the excellent oxidation resistance from synergistic effects [12], A protective Cr203 scale is formed during high-temperature oxidation, which prevents the rapid oxidation of iron and oxygen dissolution into the alloy. Since the transient Cr2o3 scale does not allow the internal oxidation of Al, Al atoms in the alloy diffuse to the alloy/scale interface and form an even slower growing Al2o3 scale. To form a Kanthal™ coating on steel by the HAPC method, a pack chemistry must be designed that will deliver Cr and Al in the desired ratio from the gas phase. Due to the large difference in stability for the volatile metal halides, this process is inherently difficult. Galmiche [73] demonstrated that a "high activity" process could result in some doping of Cr into an aluminide coating on a Ni-base superalloy. Millet [74] developed a chromizing and aluminizing codeposition process for Ni-base superalloys, which is actually a two-step process because the Al is deposited first and followed by Cr deposition. Codeposition of Cr and Al in a single step has been achieved experimentally by using a Cr-Al masteralloy [75-77]. Chromium-siliconizing in a single step was achieved by Harper [78] through the use of a mixed activator with a reactive Si02 filler. Mueller et al. [25-27] showed that a Ge-doped MoSi2 coating results from using a mixture of pure Si and Ge powders with the proper combination of activator and filler. 19
Dzyadykevich and Zablotskaya [55] describe the codeposition of B and Si in a single step, but the coating is a Si-doped TiB2 layer with inferior oxidation resistance. Three Russian patents have been issued on the codeposition of Si and Cr on titanium and other refractory metals using the HAPC method [79-81]. Since the low thermodynamic stability of chromia is a serious issue, the coatings achieved with Cr and Si are not very protective. Ho other information is available in literature on the codeposition of Si and B or Si and Ge on Ti in a single step by the HAPC method. 2.2 Thermodynamics and Kinetics of the Pack Cementation
Process
2.2.1 Single Element Deposition The thermodynamics and kinetics for depositing a single element by the HAPC method have been investigated and discussed by many authors [63-66,82-90]. Levine and Caves [66] conducted a systematic kinetic study on the effect of activator, time and temperature for the growth of aluminide coatings on a Ni-base superalloy. The models derived by Levine and Caves [66] have been found to be applicable to the growth of most aluminide coatings. The kinetic process for the HAPC method can be divided into 4 reaction steps: (1) generation of volatile metal halides in the powder pack, (2) gas-phase diffusion of the metal halides to the surface of the workpiece, (3) deposition of metal from the halide molecules at the surface of the workpiece, and (4) solid-state diffusion of the coating elements into the workpiece [83]. During the isothermal coating anneal, the halide activator dissociates and reacts with the masteralloy to form volatile metal halides by Eq.
(1.2). Equilibrium between the activator, masteralloy and filler is quickly established within the pack and maintained throughout the coating treatment [82,83]. Computer-assisted thermodynamic calculations are an accurate prediction for the vapor pressures of the metal halides during the coating treatment [75-78,82,83]. The deposition reactions at the surface of the workpiece generally occur at a fast rate, and local equilibrium is assumed at the gas/coating interface. Therefore, gas-phase or solid-state diffusion are usually the rate-limiting step for the HAPC method. A Fick's first law expression can be used to model the gas-phase flux (Jg) in the pack, Jg o DitdCi/dx) (2.1) where is the gas-phase interdiffusion coefficient of a metal halide, cL is the concentration (c^n^/V-Pj^/RT) and x is the distance. The gas-phase flux is described by diffusion through a depleted zone of width S adjacent to the workpiece, which is commonly observed [66,84-86],
j - r ~PiAPi * ** *RT (2.2)
Levine and Caves [66] derived an expression for the steady- 21
state deposition rate (dw/dt In g/cm2-sec) from Eq. (2.2), which accounted for the porosity (e) and tortuosity (1) for
gas-phase diffusion in the pack, . (2.3) = (£.) i eMiJob) = Jk dt ' w* ' i ' w
where p is the density, the molecular weight of the element being deposited, and the tortuosity equals the reciprocal of porosity for a porous powder medium. Integration of Eq. (2.3) yields the parabolic gravimetric rate constant (kg) for the coating weight-gain resulting from gas-phase diffusion controlled growth,
kg = 2pe2Jg$ (2.4) The weight gain kinetics that are predicted by Eq. (2.4) were a close match to the measured kinetics for packs that used
either a more stable or non-condensed activator [66]. This result indicates that gas-phase diffusion through the depletion zone was the rate-controlling step for these activators. Low vapor pressures for the metal halides result
from using a more stable activator, and Eq. (2.2) shows that the slow gas-phase flux produced gas-phase controlled kinetics. Venting of the volatile, non-condensed activators (NH4X: X**C1, Br, I) was thought to limit the kinetics of coating growth by gas-phase diffusion [66,83], However, solid-state diffusion was shown to be dominant for coatings grown by packs that consist of less stable activators. 22
Levine and Caves [66] defined two mechanisms for deposition from the gas phase that are based upon the condensed state of the activator during the coating treatment, as shown in Fig. 3 for the activator circulation or condensation mechanism. Volatile activators, which do not form a condensed phase at the coating temperature, deposit the masteralloy elements by the circulation mechanism and the disproportionation reaction (Eg. 1.5). Figure 3a shows that the higher metal halide recirculates back into the pack through a depleted zone to form a lower metal halide and repeat the deposition process. The condensation mechanism is observed when the halide salt is a solid or liquid phase at the coating temperature. Figure 3b illustrates that the higher halides do not recirculate back into the pack, but the salt condenses on the surface on the coating by Eq. (1.6). Kandasamy et al. [85] modified the condensed activator mechanism by describing two overlapping depleted zones; (1) metal halide depleted zone and (2) halide plus masteralloy depleted zone (Figure 4a). Kandasamy et al,[85] described the aluminizing of Ni via an A1F3 activator by characterizing how much activator in pack results in the circulation, condensation or mixed circulation~condensation mechanism, as summarized in Fig. 4b. An adherent layer of activator salt is always observed on the surface of coatings grown by the condensation mechanism. Furthermore, Kandasamy et al. [85] showed in Fig. 4b that the growth kinetics must be independent 23
UNDEPLETED PACK DEPLETED SPECIMEN ^XWSXWXXXXXXXXXXWXXXXXXX k\XV\XX\\XXV\AX\\XS\XXX\\X\ vX\\\\\XVSV\\NN\\\\\\S\X\XX AX\XX\\XXXXX\\X\\X\\X\\X\X hA\\\XX\\X\\X\\XX\\X\\XXXXXX\X\\\XV\\X\\\VS\\\XX\\\NS ASXVXXX\\X\\\\\\\\\VX\A\\X A\\V\X\\\V\SV\\XX\XXV\xXXXAWWWWWWXXXWXNWXNX v\N\X\XXSS\\\\\\\XX\\SXX\\X AXNXW\XXXXXV\\\XXX\\VX\\\A\\V\X\V\\V\\\\SXX\\S>X\W X\\X\\XX\\\\\\\XX \\X\XX\\\ X \\\\\\\X \\\\\\\\\\X \\X \\\ \\\\\\>\\X\X\SXXX\\\\NX\VS AXXXVXXWXXXXXXXXXXXXXXXXX k\\X\\\\AXX\\\X\XX\\\\\VXX\ XXX\\\XXX\\\\%\X\XX\XXXX\X k AI \\\\X\\XVX\\\ ■ 11 11 W Al xxxxxxxxxxxxx AX\X\\XX\X\\VX\XX\\\AX\X\\ >AXX\\X\X\X\X\\\X\\\\XXX\\X AXXX\\X\XX\X\XXXXXX\XX\X\\ rvx\\\xx\x\x\xx\\x\\\xxx\\\ AXNVWWXWXWXWWSVVVWV |AX\VXX\\X\XX\\X\VXX\XXX\\\ xxxxxxxxxxxxxxxxxxxxxxxxxv xx\x\\\\xxx\xxxxxx\\\xx\xx XAXX\XX\X\XX\\XXXXXX\\XXXX X\\XXX\XX\\\XXXXX\\\\XXXXX \X\\XK\X\X\X\\X\X\X\XX\\X\ X\X\X\\X\XXXXXX\XXXX\\XXX\ XX\\\XXXXXXXXXXXXXX\\XXX\\ ,x\\xx\\xx\xx\xx\xxx\x\\xxx XXXXXXXXXXXXXXXXXXXXXXXXXX XXX\\XX\X\\K\XXX\\X\\\X\XX \\\X\\W\\SSN\\XN\\NN\\\\N .XXXXXXXXXXXXXXXXXXXXXXXXXX (a)
XXX\XX\\\XX\XX\\X\XXXXX\XXX XXXXX\X\X\X\\X\X\X\\X\X\XX .XXX\X\\X\XX\XXXXXXX\\XXXXX AXXXXXXXXXXXXXXXXXXXXXXXXX A\VAXXXN\XX\\X\\XXX\\XXXX AXXXXXXXXXXXXXXXXXXXXXXXXX AXXXXXXXXXXXXXXXXXXXXXXXXX AXXXXXXXXXXXXXXXXXXXXXXXXX A\XXW\V\X\\NV\\X\XX\V\XXXAXXXXXXXXXXXXXXXXXXXXXXXXX AXXXXXXXXXXXXXXXXXXXXXXXXX Akxxxxxxxxxxxxxxxxxxxxxxxx \XXXXXX\XXXX\XXXXXXXXXXXX\XV\XXXSXX\NXN\\\VXXV> J^ k A* l1 XXXXXXXXXXXXXXWWXWWWXXX AWWWWNVXXWVVSWWXXXX^ H l XX\X\X\\XXXX\XXXXXXXXXXXXXXX \\N\\WV\\XX\XS\\N\S>SXSXXAXXXXXXXXXXXXXXXXXXXXXXXXX AXXXXXXXXXXXXXXXXXXXXXXXXX KxXXXXXXXXXXXXXXXXXXXXXXXXX AXXXXXXXXXXXXXXXXXXXXXXXXX
AXXXXXXXXXXXXXXXXXXXXXXXXXXX\\V\\\V\\\\X\VX\XNX\X>'' AXXXXXXXXXXXXXXXXXXXXXXXXX \x\\xxxxxxxx\xxxx\\x\\xxxx xx \\ xxxx \ x x \ xxxx \\\\\ x \ x a XXXXXXXXXXXXXXXXXXXXXXXXXX XXXXXXXXXXXXXXXXXXXXXXXXXX AXXXXXXXXXXXXXXXXXXXXXXXXX AXXXXXXXXXXXXXXXXXXXXXXXXX XXXXXXXXXXXXXXXXXXXXXXXXXX (b)
Figure 3. Schematic illustration of the (a) circulation mechanism, and (b) condensation mechanism described by Levine and Caves [66] where circles are Al(l) and the triangles are A1F3(c). 24
ACTIVATOR ACTIVATOR+AI UNDEPLETED PACK DEPLETED DEPLETED SPECIMEN ZONE ZONE
XXXXXXXXXXXXXXXXXXXXXXXXXX XAAXXAXXXXXXXXXXXXXXXXXXAA xxxxxxxaxxxxxxxxxxxxxxxxax xxxxxxxxxxxxxxxxxxxxxxxxxa XXAXAXaXAAAXWAXAXXAXAAXAA AAAAA\XAA\AAAAA\A\A\\\AA\X
XAXXXAXXXXxAXXAXXAXAXAAAXX AXAAAWAXAAXWXXAAXXAXXXXX AAAAA\AA\XA\\AA\XAXA\XA\A\ XAAXAXAAXXAXAXAAAAAXXAXXAA \\XAAA\AA\\X\\AA aAAXA%AAAA XXXXXXXXXXXXXXXXXXXXXXXXXXxxaaxaaaaaxxaaaaaaaxaaxxaa xxxxxxxxxxxxxxxxxxwxxxxxx XXXXXXXXXXXXXXXXXXXXXXXXXX XXXXXXXXXXXXXXXXXXXXXXXXXX XXXXXXXXXXXXXXXXXXXXXXXXXX XXXXXXXXXXXXXXXXXXXXXXXXXX AXXXXXXXXXXXXXXXXXXXXXXXXX *- Al XXXXXXXX " ' w Al XXXXXXXX UMitiitfM M *«xxxxxxax Xxxxxxxxxxxxxxxxxxxxvxxxxx xxxxxxxxxwxxxxxxxxxxxxxxx AXXXXXXXXXXXXXXXXXXXXXXXXX 4 . AXXXXXXXXXXXXXXXXXXXXXXXXX t AXXXXXAxXXXXXXxXAAAXAAAXAX - AXXXXXXXXXXXXXXXXXXXXXXXXX AXXXXXXXXXXXXXXXXXXXXXXXXX AXXXXXXXXXXXXXXXXXXXXXXXXX AIF.-* a AXXXXXXXXXXXXXXXXXXXXXXXXX AXXXXXXXXXXXXXXXXXXXXXXXXX xxxxxxxxxxxxxvxxxxxxxxxxxx Sxxxxxxxxxxxxxxxxxxxxxxxax XXXXXXXXXXXXXAXXXXXXAAXXX xxxxxxxxxxxxxxxxxxxxxxxxx Kxxxxxxxxxxxxxxxxxxxxxxxxx XXXXXXXXXXXXXXXXXXXXXXXXXX AXXXXXXXXXXXXXXXXXXXXXXXXX axxxxxxxxxxxxxxxxxxxxxxxxx aX xxxxaaaaxaaaxxx*\ x* xaxax axxxxxxxxxxxxxxxxxxxxxxxxx XXXXXXXXXXXXXXXXXXXXXXXXXX AXXXXXXXXXXXXXXXXXXXXXXXXX AXXXXXXXXXXXXXXXXXXXXXXXXX ?
ttufl MCCnAMftM' CO«NN|llM
Figure 4. Extension of the condensation mechanism, (a) Mixed mechanism of Al transport for a condensed activator with an activator and activator/source depleted zone, (b) three regions described for the condensed activator mechanism for the aluminizing of Ni by an A1F3 activator [85]. 25 of activator content for the condensed activator mechanism. The vapor pressures for the metal halides in the pack are decided by the thermodynamic stability of the halide activator. A more stable activator results in lower vapor pressures, while a less stable activator results in higher vapor pressures. Since the gas-phase flux that results from using a more stable activator is low (Eg. 2.2), gas-phase diffusion through a depleted zone may control the growth kinetics for the coating. In principle, solid-state diffusion in the coating layer should become dominant with a higher gas- phase flux. If solid-state diffusion is truly the rate- limiting step, then the growth rate should be independent of the specific activator used in the pack. Previous work on aluminizing pure Ni [84-86], pure Fe [83,87], Ni-base superalloys [66] and a Ti-Al-Nb alloy [89] have shown that faster deposition rates and thicker coatings result from using a less stable activator. The Al concentration at the surface was increased by using a less stable activator [83-88]. This trend was observed for both the circulation and condensed activator mechanism. The gas-phase diffusion model of Levine and Caves [66] provides an excellent prediction for the deposition rates of aluminide coatings (Eg. 2.4) [83-89]. Thus, the gas-phase diffusion step has the greatest influence for pack aluminizing.
The kinetics for the chromizing of iron are complicated by a displacement reaction between the Cr halides and the iron 26
workpiece: CrX2(g) + Fe(s) - FeX2(g) + Cr(s) (2.5) The deposition reaction was changed to hydrogen reduction by the addition of hydrogen to the inert gas [68,90,91], CrX2(g) + H2 (g) - 2HX(g) + Cr(s) (2.6) The displacement reaction is most important for chloride-base activators, but the high CrF2 and low FeF2 vapor pressures that result from using a fluoride activator are thought to suppress the displacement reaction [68]. O'Connel [92] chromized Fe-Cr alloys in Ar using less stable FeCl2 and CrCl2 activators and a more stable NaCl activator. The displacement reaction dominated the deposition of Cr for the less stable CrCl2 and FeCl2 activators, and the same Cr content was deposited at the surface of the workpiece for each activator. The rate-limiting step for the growth of the chromized coating was solid-state diffusion for the CrCl2 and FeCl2 activators. The deposition of Cr by the more stable NaCl activator was controlled by gas-phase diffusion, and the displacement reaction was not observed. Mazille [93] chromized Ni and observed that the kinetics were controlled by solid-state diffusion. In general, solid-state diffusion is the rate-limiting step for chromizing by a less stable activator.
Little work has been done to study the fundamental mechanisms for the kinetics of pack siliconizing. Guille et al. [54] siliconized Ti, and observed no change in the growth kinetics for a more stable activator. Mueller et al. [25-27] 27 and Page and Bartlett [94] siliconized Mo, and showed that the rate-limiting step was solid-state diffusion. Similar to chromizing, these results indicate that siliconizing is generally controlled by solid-state diffusion. Other variables that influence the deposition kinetics are the filler, inert gas and activity of the masteralloy. Kung and Rapp [82] used computer-assisted thermodynamic calculations to determine the metal halide vapor pressures for aluminizing, chromizing and siliconizing in Ar, H2/Ar and H2 atmospheres. The addition of hydrogen resulted in a significant decrease in the metal chloride or fluoride pressure for packs containing either Si or Cr by the formation of HCl of HF, but little change was observed for the Al chlorides or fluorides. The activity of the source powder is generally lowered by using a masteralloy; e.g. a Ni-Al alloy is used for aluminizing instead of pure Al. The most prominent change in the deposition kinetics that result from using a masteralloy is a decrease in halide vapor pressures. A lower activity source also changes the chemical potential gradient in the pack, and the activity for the coating layer on the workpiece cannot exceed the activity of the masteralloy. As previously discussed for a "low activity" aluminizing process, NiAl is deposited instead of Ni2Al3, and the growth mechanism is changed from the inward diffusion of Al to the outward diffusion of Ni. 28 diffusion of Ni. If the filler is not inert but reacts with the activator, then the magnitude of the metal halides can be changed. Harper [78] showed that the reaction between the Si02 filler and NaF activator to form sodium silicate increased the vapor pressures for the metal fluorides.
2*2*2 Codeposition The codeposition of two or more elements in a single processing/reaction step by the HAPC method requires an equivalent gas-phase flux for the metal halides [35]. Equation (2.2) indicates that the gas phase flux is largely determined by the partial pressure for the metal halides. Thus, the processing conditions for codeposition are achieved by devising a pack chemistry that produces similar vapor pressures for the metal halides of two elements. Figure 5 shows the partial pressures of metal halides for a pack containing a NaCl activator with Al, Cr and si. The thermodynamics for this system are such that the Al chlorides dominate the Cr and Si chlorides, and an aluminide coating with no Cr or Si additions should form by this pack. The pack composition must be modified to achieve codeposition of either Cr or Si with Al. The most successful method for achieving codeposition is to use a binary alloy powder as the masteralloy. The Cr-Al system exhibits large negative deviations from ideal behavior, and Ravi et al. [75,94] determined the masteralloy composition 29
ajci
i qT s
s o
10 it it.9 12 t3 u IS Twnpwitunxio^pq
Figure 5. Log partial pressure as a function of temperature calculated for a pack containing a NaCl activator and pure Al, Cr and Si [35]. and the halide activator required for codeposition by computer-assisted thermodynamic calculations. The codeposition of Cr and Al was achieved on low alloy steels [96], stainless steels [97] and Ni-base superalloys [35,75-77] by using a Cr-Al masteralloy with a chloride activator. Bianco and Rapp [76,77] showed that Cr-Al can be codeposited with a third element (reactive element (RE)) on Ni-base superalloys by using a pack comprised of a Cr-Al masteralloy with either a RE-chloride activator or a chloride activator with a RE-oxide filler. A NiAl coating layer with cr-Cr particles and 0.1 to 0.2 at.% RE was produced by this pack.
The growth kinetics for the Cr-modified NiAl coatings on pure Ni that result from Cr-Al codeposition were determined by Ravi [75] to be slightly slower than the growth rate of a pure NiAl layer [86]. Bianco and Rapp [76,77] measured the growth kinetics for the Cr-modified NiAl coating with RE additions on a Ni-base superalloy by two pack arrangements: (1) the workpiece in direct contact with the pack powder, and (2) the
"above pack" technique where a porous filter was used to separate the pack powder from the workpiece. The kinetic measurements for the direct contact arrangement were tainted by the inclusion of pack particles into the coating, which results from the outward diffusional growth mechanism for the "low activity" process. However, the growth kinetics were comparable to the rates measured by Tu and Seigle [98] for aluminizing Ni-Cr alloys, and the activation energies were in agreement with reported values for solid-state interdiffusion in NiAl. These results demonstrate that: (1) the growth kinetics for a Cr-modified NiAl coating, in which Cr has limited solubility in NiAl, are not significantly different than the aluminizing of Ni-base superalloys, (2) solid-state diffusion is the rate-limiting step. Bianco and Rapp [76,77] observed that the growth kinetics and activation energies were much lower for the "above pack" arrangement at the same temperatures. The porous sieve used to provide the "above pack" arrangement is a barrier to gas-phase transport and changes the rate-limiting step from solid-state diffusion to gas-phase diffusion. Levine and Caves [66] also observed a lower activation energy for the growth of aluminide coatings by more stable activators, which were controlled by gas-phase diffusion. Harper and Rapp [78,99-102] used computer-assisted thermodynamic calculations to show the following effects of the activator on packs containing a Cr-rich Cr-Si masteralloy; (1) a chloride activator results in large amounts of Cr deposition with little or no Si, and (2) a fluoride activator deposits more Si and less Cr. These observations led to the idea of using a "dual activator" mixture of fluoride and chloride salt with the Cr-Si masteralloy to produce the codeposition of Cr and Si in the desired ratio. The sio2 filler was shown to react with Na from the NaF and NaCl activator to form Na2Si205, which increases the fluorine and 32
chlorine activities in the pack. This process was used to coat low alloy steels, ferritic and austenitic stainless steels, and resulted in a composition of 27 to 24 wt.% Cr and 5 to 3 wt.% Si at the surface. The superior performance of this coating was demonstrated for high-temperature isothermal and cyclic oxidation, oxidizing/sulfidizing, erosion-corrosion and room temperature aqueous corrosion testing. Geib [35,36] showed that Fe3Al can be aluminized to form a protective FeAl layer and demonstrated that the codeposition of Al and B is possible. A pack consisting of a Fe-Al masteralloy with a FeB compound and a NaF activator produced a FeAl layer that contained 1-3 at.% B. The purpose of the boron addition was to modify the brittle nature of the FeAl
compound. The average Vickers microhardness number was reduced from 500 to 350 for the B-doped FeAl. Mueller et al. [25-27] used an unalloyed mixture of Si and Ge powders to achieve the codeposition. The phase diagrams for the Mo-Si and Mo-Ge systems in Figs. 6 and 7 show that the same intermetallic compounds are present in each system, which indicates that mutual solubility should exist for the silicide and germanide compounds. Computer-assisted thermodynamic calculations were used to determine the correct activator and inert filler to produce the Mo(Si,Ge)2 or (Mo,W) (Si,Ge) j compounds with 2-7 at.% addition of Ge. The lowest Ge content was observed at the outer surface of the
Mo(Si,Ge)2 coating with the maximum Ge amount at the Weight Percent Silicon 40 woo-
40 E noO' J (Mo)
(SI) 1(00 4 0 SO 00 ■0 Mo Atomic Percent Silicon
Figure 6. Phase diagram of the Mo-Si system [57].
Weight Percent Molybdenum too
MOO
(Mo)
noo
so to oo ao to too Ge Atomic Percent Molybdenum Mo
Figure 7. Phase diagram for the Mo-Ge system [57]. 34
Mo(Si,Ge)2/Mos(Si,Ge)3 interface. The Ge gradient was explained in terms of chemical demixing in a chemical potential gradient [103]. Mueller et al. [25-27] also showed that the amount of Ge in Mo(Si,Ge)2 could be increased by raising the ratio of Ge to Si in the pack. Figure 8 compares the cyclic oxidation resistance of the undoped and Ge-doped MoSi2 coating on Mo, which clearly demonstrates the improvement that results from the Ge addition. Cyclic oxidation of the Ge-doped silicide coatings on Nb alloys at very high temperature (1370 and 1550BC) showed these coatings were extremely protective. The Ge profile observed in the as-coated substrates is reversed after
M0 S12 + 7X Ge on Mo
MoSU on Mo
- 2 0 -
-30 0 5 10 15 20 2 5 i of 1 Hour Cycles ot 1370°C
Figure 8. Cyclic oxidation kinetics at 1370°C of undoped MoSi2 coatings compared with Ge-doped MoSi2 diffusion coatings on pure Mo [25]. 35
temperature oxidation: i.e., the Ge maximum is at the (Si,Ge)02/Mo(Si,Ge)2 interface. The Ge profile observed for the as-coated Mo(Si,Ge)2, which results from chemical demixing, and the reversal of the Ge profile after oxidation indicate that the solid-state diffusion of Ge in Mo(Si,Ge)2 is faster than Si. Mueller et al. [25] analyzed the kinetics for the growth of MoSi2 by equating the solid-state flux (Js) with the gas- phase flux (Jg),
* V dt where dx/dt is the change in MoSi2 thickness as a function of time and V is the molar volume. The growth of MoSi2 was parabolic, and the parabolic rate constant (kp) was measured, dx/dt - kp/x (2.7) The gas-phase flux for the Si fluorides was determined by Eq. (2.2). The gas-phase diffusion constant and depletion zone width (6) were calculated by making reasonable assumptions for the vapor molecules. Computer-assisted thermodynamic calculations were used to determine the partial pressures of
Si fluorides in the pack and at the MoSi2/gas interface, which is the partial pressure gradient in Eq. (2.2). The activity of Si at the MoSi2/gas interface waB determined by equating the gas-phase flux with the solid-state flux. Since the Si activity gradient in the solid phase (Aasis0.821) was much larger than the gas-phase (Aasi=0.17), solid-state diffusion 36
is the rate-limiting step for the growth of the MoSi2 coatings. Additionally, the activation energy for the growth of the Ge-doped MoSi2 coatings agrees with the values reported for Si diffusion in MoSi2.
2.2.3 Growth Kinetics Gas-phase diffusion controlled kinetics for coating growth were described by Eqs. (2.2) to (2.4), and solid-state diffusion controlled growth is the other step to consider. Two limiting cases of coating growth are considered: (1) an alloy phase where complete solubility exists between the workpiece and coating elements that are deposited at the surface, and (2) an alloy phase where little solubility exists between the workpiece and coating elements, so that deposition results in the growth of a compound with a narrow range of stoichiometry. Fick's second law is solved to predict the concentration profile for an alloy system [104,105],
<2*8> S i where c is concentration, t time, x distance and D the diffusion coefficient. Equation (2.8) can be solved for a binary alloy by assuming that a constant surface concentration is maintained with the gas phase and that the workpiece is a semi-infinite solid,
C_C^- = erf {— — ) c„-c.*o 37 where ca is the surface concentration, cQ is the concentration at all x at t“0, and c is the concentration at x and tine t. The solution of Eq. (2.8) is nore conplex for a ternary alloy A-B-C, where B and C are the inward diffusing elements,
dCfl QDbb &cB ±nA J dDgc dcc ' nA & c c dt dx dx BB dx2 dx 30 dx2 (2.10)
dcc _ dD$c dcCnA d>cc . dD& dcB nA d^Cg “dt ^ ^ D
dD m dD dD dcc ~5x “ I c ^ H x !)Fc ! b c (2.12)
Ravi [75] showed that a finite element approach can be used to solve Eqs. (2.10) and (2.11),
j A L - t < £ 0
"jr^ a t
(c/a )c = {cc) ^ + - ^ ^ n c c)^1-2(c(7)^+{cc)^1] + (Ax)'
D ^ t I (cb)^*i"2 (cb)%+ (cfl) (2.14) (Ax)3 where c^ is the concentration at the center of the ith element at the jth instant of time. The boundary conditions used to solve the problem were: (1) the workpiece was assumed to be semi-infinite, (2) a constant surface composition was maintained by the gas phase, and (3) B or C were not initially present in the workpiece. The use of small elemental grid size is required for an accurate solution of any finite element problem [105], and the time increment chosen by Ravi [75] was such that,
(2-15> At s 0.2 D
Excellent agreement was observed between the predicted and the experimentally measured diffusion profiles. The second limiting coating model involves the diffusive conversion of the substrate into a compound that has a narrow 39
range in stoichiometry. If the growth of this phase is * controlled by solid-state diffusion, then the Wagner theory should predict the experimentally measured parabolic rate constant (Eg. 2.7) for the growth of a thin, compact, single layered compound [105-107]. Equilibrium is assumed at the coating/workpiece and coating/gas interfaces. The chemical potential gradient across the compound provides the driving force for the quasi-steady-state diffusional growth of the coating. The Wagner model allows for the simultaneous transport of electrons, charged cation and anion vacancies. However, the vacancies are not required to be charged, so that the theory is valid for an intermetallic compound that is a
metallic conductor. The rational parabolic rate constant (kc in equiv. per cm-s) for the growth of an electronic conducting, single coating layer is predicted by the Wagner theory as [105-107]:
(2.16)
where Coq is the number of equivalents/cm3 for the MaXb layer, ax is the activity of the element forming the coating, and is the self-diffusion coefficient. The practical rate constant (kp, in [cm2/s]) is defined in terms of the layer thickness (£)# £2 = 2kpt (2.17) which is related to the rational rate constant [105-107]: 40
kp « kr/ceq (2.18)
and Ce(J is defined as,
c m z xP ti,si (2.19)
where p is the density, Zx is the charge of the anion, and M is the molecular weight. From Eqs. (2.16), (2.18) and (2.19)
and the further assumption that the diffusion constant is independent of composition, which is reasonable for diffusion in a line compound, then for the growth of a single layer, kp ■ I D(dlnax) = Df (dlnax) (2.20) Thermodynamic equilibrium is assumed at the M/MVX and MvX/gas
interfaces for layer growth: / dlnax ** ln(a''x/a'x) ■ In 1 - In a'x = -AG°MvX/RT (2.21) where R is the gas constant, T the absolute temperature and AG°MvX the standard Gibbs energy of formation per mole of X, which gives the following theoretical expression for the growth of a single layer:
k p - - d (A g ° M6VX/ r t ) (2 .22) Rapp [107] reviewed the literature and showed the Wagner expression provides an excellent prediction of the experimental rate constants for a variety of cation-diffusing scales on metals.
2.2.4 Growth of Multi-layered Coatings The theoretical treatment for the diffusion controlled growth of multi-layered compounds is more complicated than the 41 single-layered Wagner theory. The simultaneous, diffusion- controlled growth of more than one layer was first treated by Kidson [108] in 1961, and was later addressed by other workers [33,109-114]. The theoretical treatment of multi-layer growth has been approached by an extension of the Boltzmann-Mantano analysis to multiphase growth [109], partitioning of the diffusional flux between adjacent layers [110,111], expressing the interface velocities in terms of a flux balance at the interphase boundaries [113,114] or accounting for solid-state reactions at the interphase boundaries [112]. These treatments of multi-layered growth produce equations that are algebraically equivalent, and are complementary because the theories are written in terms of different experimental variables. The measured growth rate for any phase in a multiphase diffusion couple is influenced by the neighboring layers. The parabolic rate constant that is experimentally measured for each layer is an apparent rate constant kp1,
d5/dt • kp '5 (2.23) which upon integration gives t2 - 2kp't (2.24) Because the actual growth rate for each layer is tainted by the presence of the neighboring layers, the apparent rate constant only describes the observable growth rate. Unlike the case of single-layered growth, a theoretical equation that is based on solid-state diffusion controlled kinetics cannot 42 be simply stated for the apparent rate constant. The intrinsic rate constant is the growth rate for a layer that is in local equilibrium with the neighboring layers at the respective interphase boundaries, without mutual interference from the adjacent layers. The problem of determining the intrinsic rate constant (kp) from the experimentally measured apparent rate constant (kp*) was first considered by Yurek et al. [110] for the growth of a dual layer oxide scale. A model was developed for the growth of a dual-layer oxide scale by considering the partitioning of flux between two layers. The assumptions of the model are quasi steady-state growth, a narrow range in stoichiometry for each compound, and diffusion controlled growth for oach layer, which gives the expression for the two-layer scale in Fig. 9a,
w ' k > ° ' 0) [1 % S ’ (2.25a)
W <» ■ * 0 * 0 ) (2.25b)
where ( and » are the thicknesses and v and S are the stoichiometries of the Mev0 and Me60 layers, and V 1 is the molar volume. Hsu [111] extended theory of Yurek et al. [110] to a three layer oxide scale shown in Fig. 9b, 43
CY Vmq,o kp (Me,0) » k'p(MevO) tl+ v«*W> + vaVH.j0 (2.26a)
a ^H0gO t CY ^W»t0 kp (MetO) - kpiMe^O) [1+ P W + flP W (2.26b)
kp (M0tO) - kp (,MetO) [1 + ( P^fa,0 . YV*«,o (2.26c) where a, B and y are the thickness and v, 6 and £ are the stoichiometry of the respective MevO, Me60 and HecO layers. Wang et al. [112] used a slightly different approach that is based upon the rate of solid-state displacement reactions or "mutual consumption" of every layer at each interphase boundary. A matrix relationship was developed for parabolic multi-layered growth involving any number of layers, R » M"1 R* (2.27) where R is a l-by-n the matrix of the intrinsic growth rates for a n-layer scale, H_1 is an n-by-n matrix of the stoichiometry (vL) and molar volume (VA) of each phase, and R* is a l-by-n matrix of apparent growth rates. The apparent rate (Rj.') of each layer is written as,
Ri' - Cl2/5l (2.28) where is the thickness of the layer, and is the apparent rate constant,
£t2 ® 2C12t (2.29) 44
The intrinsic growth rate (R^) of each layer is, Ri “ ki/5* (2.30) where kL is the intrinsic rate constant. Substitution of Egs.
(2.2B) to (2.30) into Eg. (2.27) gives, (kt/Ci) « M_l CL (2.31) Wang et al. [112] showed that the sane equations given by Yurek et al. [110] (Egs. (2.25a) and (2.25b)) and Hsu [111]
(Eqs. (2.26a) to (2.26c)) are derived by Eg. (2.31). The approach derived by Wang et al. [112] does not consider the charge of the diffusing species, and is independent of a inward or outward diffusion nechanisn. The intrinsic rate constant for any compound (MevX) in a n-layer series can be theoretically predicted by using a modified form of the Wagner theory [110,111], (2.32) kp(MevX) ■ (D) [dln(ax)l ** (D) fdln(ax) where D is the interdiffusion coefficient and ax is the thermodynamic activity of the MevX compound. If the growth of each layer is controlled by solid-state diffusion, then the intrinsic rate predicted by Eq. (2.32) should be equivalent to the experimentally determined intrinsic rate constant that is calculated from the apparent rate constant. 2.3 Performance of Coatings Three criteria can be used to judge the performance of a protective coating during high-temperature oxidation: (1) rate of coating interdiffusion with the substrate, (2) the 45
M«„0 MejO Ma Ojtg) t*- •• -V'**»« - ■ V*' ■
P0|(Mt/Ma,0) P0|(MerO/Mt^) po. (a)
PO,^U/M«vP) Pot^U*p/H*|0) Po^aOJ/H-dO) JjO
Ot(|)
(b)
Figure 9. Schematic for the (a) dual-layer oxide scale described by Yurek, Hirth and Rapp [110], and (b) the triple layer oxide scale described by Hsu [ill]* 46 oxidation kinetics of the coating (both isothermal and cyclic oxidation), and (3) detection of substrate contamination to determine whether the coating was an effective barrier against the inward penetration of oxygen. Since the purpose of coating a workpiece is to extend the usable life of the component at high temperature, a mechanical test of a coated part that simulates the working environment is also required.
2*3*1 Interdiffusion For the case of coatings that form slow growing A1203 or Si02 scales during high-temperature oxidation, the coating lifetime is generally not limited by the diffusive loss of the coating to form the oxide, but rather interdiffusion of the coating into the substrate. Interdiffusion is a serious issue when the lower compound of the coating phase has poor oxidation resistance. For example, the life of MoSi2 coatings on Mo alloys is limited by the rate at which Mo5Si3 consumes MoSi2 upon the dissolution of Si into the Mo substrate [10,94,115-119]. Since the high-temperature oxidation resistance of Mo5Si3 is very poor, the consumption of the outer MoSi2 layer by Mo5Si3 limits the life of these coatings. However, Mueller et al. [25-27] demonstrated that the dissolution rate of a (Mo,W) (Si,Ge)2 coating is much slower. Smialek and Lowell [120] have shown that interdiffusion of NiAl coatings with Ni-base superalloys can be life limiting. Thermodynamic calculations have shown that a slow growing Si02 scale is stable on every Ti silicide in the Ti-Si system 47
with the exception of Ti3Si. Therefore, interdiffusion between the Ti substrate and its silicide coating would not be
expected to limit the life of the silicide coating.
2.3.2 Isothermal and Cyclic oxidation As previously mentioned, the growth rates for Si02 and A1203 scales are so slow that the scaling rates for compounds or alloys that form these oxides are not usually an issue. One notable problem for compounds that form Si02 or A1203 scales is an accelerated rate of oxidation of the more mobile
component at low temperature [31]. One of the best known examples is MoSi2, which has excellent oxidation resistance at high-temperature but can undergo rapid, destructive oxidation at low temperature (posting). The posting of MoSi2 will be addressed in the proceeding section. The slow oxidation kinetics for TiSi2 coatings and thin films at temperatures above 700°C results from the growth of a continuous Si02 scale [52,56,121-122]. A faster rate for the oxidation for TiSi2 is generally observed at temperatures below 700°C due to the formation of a mixed Si02-Ti02 scale [52,121-122]. The low-temperature oxidation of TiSi2 is not catastrophic, but the rates exceed those extrapolated from high temperature. Ti02 is not volatile and the growth rate for Ti02 is rather slow at temperatures below 700°C [10].
Thus, the low temperature oxidation of Tisi2 is not destructive nor rapid enough to be a major concern [121-122]. However, Meier et al. [31] has reported that Ti5Si3 is 48 susceptible to a form of pesting at low temperature; this contradicts reports by other authors [21]. Most parts used in high-temperature applications undergo a periodic cycle of heating and cooling in service. Therefore, the protective coating must withstand many thermal cycles in an oxidizing atmosphere. The best way to evaluate the performance of a coating is cyclic oxidation. A large mismatch in the CTE between the coating and oxide scale, which is brittle at room temperature, may cause spelling during the cooling cycle. Nesbit [123] and Lowell [124] have developed models to predict the lifetime of NiAl coatings that are based upon spalling of the protective a-Al203 scale. The repeated loss and formation of a-Al2o3 can lower the A1 content to the point where the coating is not protective. A major concern regarding the use of brittle compounds as protective coatings is cracking upon thermal cycling that results in oxygen contamination of the workpiece. Smialek et al. [89,125] used an aluminizing treatment forming a TiAl3 compound to protect a Ti-24Al-llNb alloy. Despite the low weight changes for cyclic oxidation, further development of these coatings was not pursued because cracking of the coating resulted in substrate contamination [14]. Therefore, while oxidation kinetics are important, a protective coating must also be an effective barrier against oxygen penetration. 49
2.3.3 Substrata Contamination
Refractory metals such as Ti, Nb, Ta and Zr have large solubilities for oxygen and other interstitial elements [10]. The dissolution of oxygen decreases the ductility of the alloy. The inward penetration of oxygen and other contaminants is generally detected by a steep gradient in hardness beneath the scale/metal interface [11,12]. Since the embrittlement of Ti alloys by oxygen dissolution is very significant, microhardness profiles are an accepted means of testing for oxygen contamination [11-15].
2.4 Pasting of HoSi2 Many studies have shown that the accelerated, low- temperature oxidation of MoSi2 ("pesting") can lead to rapid disintegration of the material in the temperature range of 300 to 600°C at a rate that is maximized at 500-550°C [29- 32,126,127]. At high temperature, Mo trioxide has sufficient volatility to result in its complete evaporation, but at low temperature the slow growth kinetics for Si02 and the low vapor pressures for the Mo oxides result in the formation of a mixed Si02-Mo0x oxidation product. The rapid growth of the voluminous Mo oxides at low temperature is thought to produce accelerated crack growth at pores, pre-existing cracks, and grain boundaries in MoSi2, which results in disintegration of the material [29-31,125]. Thus, the formation of Mo oxides at low temperature is the primary cause for the destructive form of low-temperature oxidation known as pesting. since all 50
MoSi2 furnace heating elements and engine components used at high temperature have cool sections at interconnects that must exist at 500°C for an extended period of time, pesting represents a serious issue. The mechanism and kinetics of pesting have been studied by many authors, but a completely satisfactory solution to the problem has not been discovered. Pre-oxidation at high temperature to form a Si02 scale that is free of Mo oxide may delay the onset of pesting. However, the stiff and brittle Si02 scale grown at high temperature may crack or suffer damage at low temperature, and this approach does not solve the pesting problem [29,126]. Meschter [30] demonstrated that the low-temperature oxidation of dense MoSi2 that contains no pores or cracks is not catastrophic, but is fast enough to raise concern about a potentially disastrous failure in a critical application such as an aircraft engine. Additionally, brittle MoSi2 is likely to crack or form pores in service, especially when it is reinforced by a second phase with a dissimilar CTE, which would accelerate the rate of pesting. Berkowitz-Mattuck et al. [128] loaded dense, initially crack-free MoSi2 in a four-point bend test at 450 to 575°C in air and observed catastrophic failure. Thus, the use of dense MoSi2 does not stop the problem of pesting. Chou and Nieh [29] allude to the fact that impurities can accelerate or decrease the rate of low-temperature oxidation for MoSi2, but no specific beneficial compositions were mentioned. Fitzer 51 and Kehr [129] showed that germanium additions improved the pesting resistance of MoSi2, but a fairly rapid rate of low- temperature oxidation was observed. To the author's knowledge, no fully effective solution for pesting has been reported in the open literature. CHAPTER III EXPERIMENTAL PROCEDURE AND MATERIALS
3.1 Mittrials Commercially pure (CP) titanium (99.9% pure, Timet,
TIMETAL 50A, Heat# T-2002) was used for most of the coating development. The CP-titanium plate was vacuum arc-melted, and worked into a 0.5 mm thick plate, and the impurity contents were below the detection limit for EDS (< 0.1 at.%). Three different classes of intermetallic-base Ti-alumimide alloys were also used in this study. The first is based on the cr2- Ti3Al compound, and has the nominal composition Ti-24Al-llNb (at.%, NASA Lewis). The Ti-24Al-llNb alloy is two phase and consists of a continuous ordered bcc phase ("beta", CsCl structure, S.G. #221) with the majority (85 vol.%) being the hexagonal a2-Ti3Al phase (S.G. #194) [8,130]. The a2-Ti3Al alloys were prepared by a powder processing method that is referred to as the "powder cloth technique", which was described in the report by Brindley et al. [131]. Mo was used for encapsulation in the hot pressing cycle and reacted with the Ti-24Al-llNb alloy to create a Ti-Mo reaction zone that was sanded off prior to coating.
52 The second class of alloy Is based upon the orthorhombic TiAl2Nb intermetallic compound. Three different orthorhombic alloys were used In this study. The Ti-22Al-27Nb (at.%) alloy consists of a continuous TiAl2Nb phase and a discontinuous
ordered bcc phase ("beta", CsCl structure, S.G. #221) [14]. The Ti-22Al-27Nb alloy was cast and forged above the beta transus temperature (Wrlght-Patterson AFB); the temperature region where the alloy Is fully beta. Two three-phase Ti- 20Al-22Nb and Ti-22Al-23Nb alloy were used In this study; these consisted of a continuous orthorhombic TlAl2Nb phase and discontinuous ordered bcc and a2-Ti3Al phases [14]. The Ti- 20Al-22Nb alloy was cast and forged at 25-50®C above the beta transus temperature (G.E. Aircraft Engines, Evendale). The average composition of the Ti-20Al-22Nb was Ti-20.7A1-22.8A1, and the alloy contained 900 ppm oxygen. The Ti-22Al-23Nb alloy was prepared from powder by hot pressing (NASA Lewis). A Y“TiAl based Ti-48Al-2Nb-2Nn alloy is the third class of material used in this study (Wrlght-Patterson AFB). This alloy was prepared by extrusion and then forged.
Commercially pure (CP) Ho foil that was 0.1 mm thick was also coated (ALFA, 99.95% pure, lot# F01C14). A MoSi2 heating element that had failed- in service was sectioned using a slow- speed diamond saw. The oxide scale was sanded off using 320 grit Sic paper prior to testing. MoSi2 that was prepared from powder (Mermann Starck, high-purity, -90 ^m) by hipping at 1700°C for four hours under 207 MPa argon pressure was also 54 studied. 3.2 Ti-fli Diffusion couple CP-titanium and a single crystal Si wafer were polished to 0.3 pm A1203, and cleaned in acetone and ethanol in preparation for the diffusion couple experiment. To avoid any oxidation or contamination, the diffusion couple was immediately loaded into a horizontal tube furnace that was purged with prepurified helium gas. The CP-titanium was sandwiched between two pieces of Si wafer, and held on an A1203 boat by a 8g weight of Nb to maintain contact. A boat of high-purity Ti turnings was positioned upstream and downstream of the couple to getter any residual oxygen. The couples were heated (18°C/min) to either 950®C, 1050®C or 1150°C (± 5°C) for periods of approximately 14, 24, 36, 48 of 60 hours.
3.3 Pack Cementation Method The same procedure for preparing the powder packs was used for coating the metals and alloys in this study. The powders were weighed, blended and mixed in a ball mill for 3-6 hours. To eliminate moisture, the mixed packs were dried at
200°C for 12 hours and stored in a desiccator. The substrates were sectioned and prepared by polishing the surface with 320 grit Sic paper, then immediately cleaned in acetone and ethanol, measured, weighed and loaded into A1203 crucibles with the powder pack. The crucibles were sealed with an Al203-based cement (Ceramabond 569, Aremco Products) and dried at 200 °C for 3 hours. The crucible packs were loaded into a resistance-heated, horizontal tube furnace for the isothermal coating treatment. Figure 10 is a schematic illustration of the experimental design used for pack cementation. A prepurified argon gas was purged through the tube during the coating cycle. The packs were preheated in the furnace at 2S0°C for 3 to 4 hours in flowing argon to ensure complete drying of the powders. Following the preheat, the packs were heated at 4°C/rain to either 950, 1050 or 1150°C for 0 to 28 hours, where a zero hour run consisted of heating the sample to the desired temperature and immediately cutting the furnace current. A 16 to 24 hour cooling period was required after the furnace current was cut. Following the coating treatment, the packs were unloaded and the coupons were cleaned in water and acetone, and then measured and weighed.
A variety of pack compositions were used to develop the best silicide coatings for titanium. Pure elemental powders of Si (-325 mesh, 99.5%, ALFA), B (-325 mesh, 99%) or Ge (-325 mesh, 99.999%, ALFA) were used as the masteralloys. Powders of the boride compounds SiB3 (99%, ALFA), CrB2 (99%, ALFA), TaB2 (99.5%, ALFA) or Ti&2 (-325 mesh, 99.5%, ALFA) were also used with Si as the masteralloy. The halide salt activators used in this work were anhydrous A1F3 (99+%, ALFA), CuF2 (99.5%, A1FA), MgF2 (-325 mesh, 99.9%, ALFA), MnF2 (99%, ALFA), NaF (-100 mesh, 99%, ALFA), or NaCl (-20 mesh, 99%, 1. Argon Gas 2. Flow Meter 3. COj Getter 4. Deslccant 5. O 2 Getter 6. Flange 7 . Fan 8. Mulllte Tube 9. Horizontal Tube Furnace 10. Furnace Controller 11. Alumina Crudble 12. Workpiece 13. Bubbler
Figure 10. Schematic of the experimental setup used coat the workpieces by the pack cementation method. 57
ALFA). The inert fillers used were Al203 (60-325 mesh, 98%,
Fisher Sci.) or sic (120 grit, Metallurgical Supply Co.). A summary of the compositions used to produce the B- and Ge- doped silicide coatings is given (in wt.%) in Tables 1 and 2, respectively. The pack compositions used to produce the boride coatings are summarized in Table 3. A preliminary evaluation for the coatings produced by these packs was obtained by measuring the weight change after 48 hours of isothermal oxidation at 1000°C in a horizontal, resistance heated tube furnace in stagnant air. The pack compositions, which were determined to be the best from the preliminary work, were used to produce coatings for further testing, as summarized in Table 4. The ratio of si to Ge in the powder pack was varied, as summarized in Table 5. The pack compositions used to produce Mo-silicide coatings on CP-Mo are summarized in Table 6. The pack composition (wt.%) 10% Si, 10% Ge, 2% NaF and SiC or 10% Si, 2% NaF and Sic was used most extensively in this work.
3.4 Solgas Calculations The assumption that equilibrium is maintained in the bulk pack during the coating process is reasonable for the majority of pack cementation processes that use a condensed activator. The calculation of the equilibrium partial pressures for the vapor species is used to predict or understand the results of a HAPC process. Since many species must be considered for the calculation of the most simple powder pack, a computer program 58 Table 1. Compositions of pack powders (in wt.%) used in the development of a B-doped silicide coating.
Pack Source Activator Filler
1. lOSi-lOB 2MgF2 a i 2o 3 2. 10Si“10B 2NaF a i 2o 3 3. 10Si-2B A1F3 AloO-j 4. 20S1-0.02B 2NaF Ai2°3 5. 10Si-10SiB3 2MgF2 a i 2o 3 6. 10Si-10SiB3 2NaF a i 2o , 7. lOSi-O.ISiB, 2MgF2 a i 2o 3 8. 7Si-6TiB, 2MgF2 A 1 & 9. 10Si-5TiBz 2MgF2 A120, 1 0 . 10Si-2.5TlB2 2MgF2 Aijo' 11. 7Si-6TiB, 2MgF2 sic 12. 7Si-6CrB2 2MgF2 A1203 13. 7Si-6TaB2 2HgF2 a i 2o 3 14. 7Si-6TiB2 2A1F3 a i 2o 3 15. 7Si-6TaB2 2A1F3 Al203 16. 7Si-6TiB2 2Cu F2 A1203 17. 7Si-6TaB2 2Cu F2 2°3 59
Table 2. Compositions of pack powders (in wt.%) used to develop the Ge-doped silicide coating.
Pack Source Activator Filler l. lOSi-iOGe 2NaF Sic 2 . lOSi-lOGo 2MgF2 Sic 3. 12Si-6Ge 2MgF2 sic 4. 12Si-6Ge 2A1F3 Sic 5. 12Si-6Ge 2CuF2 sic 6. 7Si-6Ge 2MgF2 A1203 7. 7Si-6Ge 2NaF ai2o3 8. 7S1-6GO 2A1F3 AI3O 1 9. 7Si-0.7Ge 2A1F3 A 1 & 1 0 . 15Si-1.5Ge 2A1F3 ai2o3 1 1 . 0Si-8Ge 2A1F3 ai2o3 1 2 . 8Si-4Ge 2A1F3 13. 8Si-2Ge 2A1F3 A 1 2 ° 3 14. 8Si-4Ge IAIF3 ai2o3 15. 8Si-4Ge 0.5A1F3 ai2o3 16. 8Si-4Ge 4A1F3 ai2o3 17. 8Sl-4Ge 8AIF3 ai2o3 18. 2Si-lGe IAIF3 Al2°3 19. 4Si-8Ge IAIF3 Al}01 2 0 . 16Si-8Ge 1a1F3 A l l o\ 2 1 . 16Si-8Ge 2A1F3 A 1 2 ° 3 2 2 . 16Si-8Ge 2CuF2 ai203
Table 3. Pack compositions used to produce boride coatings on CP-titanium (wt.%).
Pack Source Activator Filler
1. 4B 2MgF2 a12°3 2. 4B 2A1F3 A1203 60
Table 4. Compositions of pack powders (wt.%) that produced the best silicide coatings on CP-titanium and were used for further evaluation by isothermal and cyclic oxidation.
Pack Source Activator Filler 1. 7Si-6TiB2 2MgF2 Al203 2. 7Si-6TiB2 2A1F3 A1203 3. 7Si-6TiB2 2CUF2 Al203 4. 7Si-6TaB2 2MgF2 A1203 5. 7Si-6TaB2 2A1F3 Al203 6. 7Si-6TaB2 2CUF2 A12°3 7. 16Si-8Ge 2A1F3 A1203 8. 16S1-8GO 2Cu F2 Al203 9. 12Si-6Ge 2HgF2 SiC
Table 5. Pack compositions (wt.%) U B e d to produce Ge-doped silicide coatings where the ratio of Si to Ge was varied in the pack.
Pack Source Activator Filler 1. 16Si-8Ge 2A1F3 A1203 2. 16Si-16Ge 2A1F3 a i 2o 3 3. 8Si-16Ge 2A1F3 a i 2o 3 4. 16Si-8Ge 2Cu F2 a i 2o 3 5. 16Si-16Ge 2Cu F2 a i 2o 3 6. 8Si-16Ge 2Cu F2 a i 2o 3 7. 12Si-6Ge 2MgF2 SiC 8. 12Si-12Ge 2MgF2 sic 9. 6Si-12Ge 2MgF2 SiC 61
Table 6. Summary of the different powder pack mixtures used to produce undoped and germanium-doped silicide coatings on CP molybdenum (wt.%).
Pack Source Activator Filler
1. lOSi 2NaF sic 2. lOSi-iOGe 2NaF sic 3. lOSi 2HgF2 Sic 4. 10Si“10Ge 2MgF2 sic 5. lOSi 2A1F3 sic 6. losi 2A1F3 a i 2o 3 7. lOSi-lOGe 2A1F3 sic 8. lOSi 2Cu F2 sic 9. lOSi 2Cu F2 a i 2o 3 10. lOSi-lOGe 2Cu F2 sic 11. lOSi MnF2 SiC was used for such thermodynamic calculations. An updated version of the SOLGAS-PV computer program called STEPSOL was used for the calculation of the equilibrium vapor pressures in a pack [132,133]. The calculation is based on conservation of mass and minimization of the energy in the system* All possible condensed and vapor species that participate in the reaction must be entered into the program for an accurate calculation. For some calculations, the reaction of the inert filler was assumed to be negligible, and was not considered. 3.5 Oxidation Evaluation 3.5,1 Isothermal Oxidation Figure n is a schematic drawing of the Cahn-171 microbalance that was used to measure the kinetics for isothermal oxidation by thermogravimetric analysis (TGA). The coated coupons were placed in an A1203 crucible that was 62
suspended from a sapphire wire, and tested at 500°C, 600“C, 700°c, soo°C, 900°c and 1000“C for 48 hours. The heating rate was 50 to 70 eC/min, and the cooling rate was 5 to 3 °C/min. Predried air from a gas cylinder was flowed at an average rate of 0.632 cm/s during each test. The weight change versus time was recorded by the Cahn DACS plus software, and plots of weight change versus square root of time were used to determine the parabolic rate constants [134]. The data were corrected for the buoyancy of the crucible by weighing the coupon before and after the test using a Mettler balance (± 0.02 mg), and subtracting the proper factor [135]. CP- titanium workpieces that were coated with a B-doped silicide (Packs #1 and #4 in Table 4) and Ge-doped silicide (Packs #7 and #9 in Table 4) were studied. The uncoated Ti-22Al-27Nb alloy, and Ti-22Al-27Nb workpieces coated with a B-doped silicide (Pack# 1) and Ge-doped silicide (Pack# 9) were also evaluated.
An open-ended, resistance-heated furnace was used for isothermal oxidation at 500°C or 1000°C for extended periods
of time. The specimens were periodically withdrawn from the furnace and weighed using a Mettler balance (± 0.1 mg). The coupons were reloaded into the furnace to continue the oxidation. 3.5.2 Cyclic Oxidation
Figure 12 is a schematic for the cyclic oxidation rigs that were built for this study. The coupons were placed into 63
16
1. Microbalance 8. Furnace switch 2. Stand 9. Status lights 3. Cooling fan 10. Elevator controls 4. Cooling Can switch 11. Thermocouple connector 5. Furnace 12. Leveling feet 6. Temperature display 13. Thermocouple connector (TG-171) 7. Gas ports 16. Vacuum port 17. Weight display
Figure 11. Schematic diagram the Cahn-171 therraogravimetric analyzer (TGA) used to monitor the weight change for isothermal oxidation. 64
a quartz holder that was lined with alumina-base refractory paper. The samples were lowered into the hot zone of the furnace and oxidized for 1 hour, and then pulled from the
furnace and cooled for 30 minutes to complete one oxidation cycle. The coupons usually cooled to low temperature (~100°C) after 5 to 10 minutes. The coupons were periodically weighed using a Mettler balance (± 0.02 mg). The duration of a typical cyclic oxidation campaign was 200 cycles. Each type of coating was oxidized at 500°C, 600°C, 700°C, 800°C, 900°C
and 1000°C, with some limited testing at 875°C. Following cyclic oxidation at every temperature, a representative coupon for each type of coating was mounted, cross-sectioned and polished using standard metallography techniques. Microhardness profiles using a Knoop indentor with a lOg load for 15 seconds were made to test for oxygen
contamination of the substrate. Several indents were taken in the middle of each coupon at a 300g load to standardize each measurement. The baseline microhardness data were established from as-coated coupons that were not oxidized. The microhardness profiles were measured from uncoated CP-
titanium, Ti-22Al-27Nb and Ti-20Al-22Nb following cyclic oxidation. A microhardness standard for the orthorhombic-alloys was supplied (Brindley [14]) late in the study which showed that the microhardness values measured at Ohio state are consistently 50-60% lower than the expected values. 65
1. D.C. Motor 2. Time Delay Relay 3. D.C Power Source 4. String 5. Position Switch 6. Quartz Tubing 7. Quartz Rod 8. Wheel 9. Mullite Tube 10. Furnace 11. Furnace Contoller 12. Quartz Holder 13. Coated Woricplece 14. Rack
Figure 12. schematic illustration of the cyclic oxidation rigs used for this study. 66
Nevertheless, the Internal comparisons and trends in microhardness values are most important for the detection of substrate contamination. Substrate contamination is revealed by a steep gradient in hardness at the surface. substrate
contamination by oxygen dissolution can also be detected by
the Kroll etch (4% HN03, 2% HF, H20 [136]) [11,12]. Microhardness was commonly used on coupons also etched to detect for substrate contamination.
3.6 Analysis, Characterisation and Evaluation of Coatings The phases at the surface of as-coated and oxidized substrates were determined by X-ray diffraction (XRD) using a SCINTAG diffractometer. Some of the coupons in the as-coated condition were sanded using 320 grit SiC paper and analyzed by XRD to determine the crystal structure of the inner coating layers. A JEOL 35A and Phillips XL-30 scanning electron microscope (SEM) was used to analyze the surface by EDS. The as-coated and oxidized coupons were mounted in bakelite, sectioned using a low-speed diamond saw and polished through 0.3 p A1203 using standard metallographic procedures. The Kroll etch was used on some coupons to distinguish each layer of the silicide coating and to study the substrate. The compositions of the coatings, oxide scales, salt layers, interdiffusion zones and substrates was determined by EDS. Elemental standards with a ZAP correction were used to quantify the results by the EDS computer software. An accelerating voltage of 20 KV was used. 67
A CAMECA microprobe with a 1 jin diameter probe size was used on other coupons to make quantitative profiles of the composition. Elemental standards were used to quantify the results. An accelerating voltage of 15 KV was used. Measurements of layer thicknesses were used to determine the interdiffusion kinetics of some coatings. The thickness of each layer was measured by optical microscopy and SEM for each combination of coating time and temperature. At least forty measurements were taken from each layer and the mean thickness was calculated for the data analysis. 3.6 Balt Coating of HoSi2 A salt layer was applied to MoSi2, the MoSi2 heating elements, undoped and Ge-doped MoSi2 diffusion coatings by spraying the coupons with an aqueous salt solution, and drying the substrates at 150°C. The amount of salt deposited was determined by weighing the coupons before and after the treatment. The aqueous solutions were prepared by heating distilled water to 85-95°C, and adding the appropriate salt until the solution was saturated. The warm, saturated solution was used to coat the coupons. NaF, NaBr, NaCl, Nal and a NaN03 mixture were used to treat MoSi2. To clarify the effects resulting from the Na-base salts, MoSi2 was coated with other substances and dried at 150°C. Tetraethylorthosilicate (Si{OC2H5)4), which forms Si02 upon heating, was used to coat some heating element coupons, other heating element coupons were coated with an aqueous BORAX 68 (Ka2B407), or an alkali silicate solution (Sermabond 487, part I, Sermatech International). Undoped and Ge-doped molybdenuro-silicide diffusion coatings on Mo, MoSi2 heating elements and hipped MoSi2 coupons that were untreated or treated with a salt layer were oxidized at 500°C for 2500 hours. The samples were periodically weighed every 200-500 hours using a Mettler balance (± 0.1 mg). CHAPTER XV
THE GROWTH AMD DEVELOPMENT OP BORON- AND GERMANIUM-DOPED
8XLXCXDE COATINGS: RESULTS AMD DISCUSSION
4.1 Titanium-Silioon Diffusion coupls Figure 13 is an optical micrograph of the etched cross- section for a solid-state titanium-silicon diffusion couple after 14 hours at 1150°C. The distinct layers of TiSi2,
Tisi, Ti5Si4, Ti5Si3 and Ti3Si are obvious upon etching, and EDS was used to verify that all five silicide phases were present in the interdiffusion zone. No oxygen contamination was detected in the Ti substrate by SEM for the diffusion couple. No hardness gradient or Ti embrittlement was detected by microhardness [11], which indicates that oxygen contamination was not present in the interdiffusion zone. The fine precipitates in the Ti substrate were determined by EDS to be Ti3Si. Since Ti has some solubility for Si at high temperature (about 2 at.% at 1150°C) [57], precipitation would be expected upon cooling. Silicon is known to be the fastest diffusing species in TiSi2 [137], and Si diffusion is assumed to be dominant in Tisi, Ti5Si4, TisSi3 and Ti3Si. Thus, the diffusional growth of the silicide layers occurs by the inward diffusion of Si through
69 70
Figure 13. Optical micrograph of the cross-section of the titanium-silicon diffusion couple following 14 hours at 1150°C. The five TiSi2/TiSi/Ti5Si4/TisSi3/Ti3Si layers are revealed by the Kroll etch.
the silicide compounds and into Ti that becomes saturated with Si.
4. l.l Theory of Five-Layered Growth The measured growth rate of any phase in a multiphase diffusion couple is affected by the growth of neighboring phases. A growing layer of TiSi in the five-layered silicide coating shown in Fig. 13 simultaneously reacts with the neighboring TiSi2 and Tissi4 layers, and the TiSi layer is in turn affected by the growth of Tisi2 and Ti5si4. This theory of "mutual consumption" as developed by Wang et al. [112] provides a tractable approach to accounting for the 71 partitioning of the flux during the multi-layered growth of compounds with narrow ranges of stoichiometry. The parabolic rate constant that is measured for each layer is an apparent rate constant kp', d*/dt = kp'* (4.1) which upon integration gives I2 - 2kp*t (4.2) The mutual consumption theory of Wang et al. [112] is directly applicable to five-layer growth by a matrix method using Eq. (2.31), (ki/CjJ - M ”1 CL (2.31) For the five-layered coating shown in Fig. 13, Eq. (2.31) results in the following:
qv 4 r iii r i ^lV3 1 q V|»1 ViVi v;vx V|*, V2v, | £ | i-S 1X Ifcj 1 Vi Viv, Viva ‘ q 1 I C . II q qv 4 (4.3) I 1 1 itei C 1 Vi V* Viv3 V ^ tel 1— 1 Va Vi X1 Vi V, v5v4 1 teJ
Vi V5 Vi i j q Vi Vi where 1 is Ti3Si, 2 is Ti5si3, 3 is Ti5si4, 4 is TiSi and 5 is Tisi2, is the molar volume, and is the stoichiometry normalized to one mole of silicon (TivSi). Matrix operation of Eq. (4.3) and substitution of Eq. (2.29) results in an equation for the intrinsic rate constant of 72 each coating layer,
kU ^-Jr'fTf k p(Ti,Si) 5i) [l+1^fit +1^ r+_r VXV454 r. ^ ^1V5^S 1 (4.4a)
k "k* (Ti Si ) r V^ 1 +1+ + ^2V4?4 ^2V5^5 i i w e t w ; w 2 3 (4.4b)
m i T . (4.4c)
(4.4d)
k -k'lTlSi ) r 1 t 1 til 5 p( 1 v,5, v,{, v,«5 v,{, 11 (4.4e) where is the intrinsic rate constant, k^1 the apparent rate constant and the measured layer thickness. The intrinsic rate constant of any layer (Tivsi) in the five-layered silicide coating is theoretically predicted by using a modified form of Wagner's theory (Eg. (2.16))
[110, 111],'
kp m vSi) i n (D ) tdln(asi) ] =-A. l (D)]t dln(asl) (4*5)
For the chemical potential gradient shown by the schematic of the five-layered coatings in Fig. 14, theoretical 73
expressions were derived for the intrinsic rate constants of
the five-layered silicide coating, kp[Ti3Si] = kp[TiSi] = (l/2)D[TiSi](lnasll5J-lnaslt4]) 4.6d) kp[TiSi2] - D[TiSi2J(-lnasi(SJ) 4.6e) where 1 is for the silicon activity at the Ti3Si/Ti interphase boundary, 2 is Ti5Si3/Ti3Si, 3 is Ti5Si4/Ti5Si3, 4 is TiSi/TisSi4, and 5 is the TiSi2/TiSi boundary. If the kinetics are controlled by solid-state diffusion, then the experimental intrinsic parabolic rate constants for each layer from Eqs. (4.4a-4.4e) should equal the theoretical rate constants in Eqs. (4.6a-4.6e), permitting a calculation of the diffusivity in each phase. Engqvist et al. [138] made a critical assessment of the thermodynamic data for the Ti-Si system, and recommended a set of data that reproduces the Ti-Si phase diagram. The Gibbs energies of formation at 1223, 1323 and 1423K are listed in Table 7, and Fig. 14 shows the activity profiles for silicon for the five-layered coating. The interdiffusion coefficient of TiSi has been measured at 275-340°C [139], and is probably a grain boundary diffusivity which does not accurately describe diffusion at high temperature. To make a general comparison, these low-temperature, grain boundary dominated diffusion data were extrapolated to 950 to 1150®C. The parabolic growth rate of a single-layered TiSi2 film that results from interdiffusion between a Ti film and a Si substrate has been measured at 475-550°c [140] and 750-950°C [141] by separate authors. Since both of these cases involved the diffusion-controlled growth of a single layer, Eg. (2.22) was used to calculate a diffusion coefficient for TiSi2. The diffusion data were extrapolated to 950-1150°C to determine the theoretical intrinsic rate constants for TiSi2 and TiSi by Eqs. (4.6d) and (4.6e); these values are given in Table 8. To the author's knowledge, diffusion data for the phases Ti5Si4, Ti5Si3 and Ti3Si are not available, so intrinsic rate constants could not be predicted. 4.1.2 Multi-Layer Growth - Global Growth Rate By extending the results of Yurek et al.[110] and Hsu [111], five equations were derived to determine the global growth rate for the entire coating from the intrinsic rate constant of each individual layer. The equation for the growth of the entire silicide coating on titanium by the partitioning of flux is (v1Pl+v2P2+v3P3+v4P4+v5P5)Ti + (P1+P2+P3+P4+P5)Si -> PjTi3Si + P2Ti5/3Si + P3Ti5/4Si + P4TiSi + P5Ti1/2Si (4.7) where is the stoichiometric fraction normalized to one mole of silicon, and PA the partitioning fraction for Ti and Si in each of the five silicide layers shown in Fig. 8. The 75 Table 7. Standard Gibbs energy of formation (AG°[J/mole]) for the five compounds in the Ti-Si system [138] at 950, 1050 and 1150°C. 1150°C 1050°C 950®C TiSi2 -172,753 -172,410 -172,155 TiSi -151,501 -152,374 -153,228 Ti5Si4 -713,884 -715,996 -718,269 Ti5Si3 -582,373 -580,252 -578,618 Ti3Si -195,454 -198,716 -201,454 Table 8. Theoretical parabolic rate constants (kp [cm2/s]) for the growth of TiSi2 and TiSi in the multi layered Tisi2/Tisi/Ti5si4/Ti5si3/Ti3si coating at 950, 1050 ana 1150°C. Diffusion data from literature are used in the calculation. 1150°C 1050°C 950#C TiSi2[140] 4.13X10“*° 1.21X10”10 2.89X10"11 TiSi2[141] 1.42X10”® 5.09X10”10 1.31X10”10 TiSi[139] 3.91X10”10 1.88X10”10 8.57X10”11 iio i h iS sse t90 15 ad15° [138]. 1150°C for and 1050 silicide profile 950, at activity system the five-layered on Ti-Si the of the in superimposed drawing is silicon that schematic A coating 14. Fig. Log Activity of Silicon - 4 - - - 2 8 - - TiSi. TiSi Ti Ti ,Si Subs rate Subs CPlr 950C- 1050(1- 76 77 total thickness (S) of the five-layered coating is given by s - 5i+52+53+^4+5s (4*®) and the global rate constant of the whole coating (K) is, dS/dt = K/S (4.9)The left-hand side of Eg. (4.9) is written in terns of each layer [111], dt dt u ^ %x %x (4.10a) dt ' d t ' 1 *2 x €2 «2 €2 (4.10b) dt dt 53 53 53 (4.10c) (4.lOd) dS dt dt (4.lOe) Hsu [111] developed equations for d^/dt for a three-layered scale by considering Eg. (4.7); extension of these results to five layers gives 78 Pi*! , kiTijSi) % dt (4.11a) P,v *L 3Ti i / frtris;33i] x dt P1V1+P2V2+P3V3+P4V4+P5V5 M 52 ' (4.iib) iii P^i jc[ri5/4si] dt PiVi+p2v2+p3v3+p4v4+p5v5 «3 (4.11c) dZi.r______P4V4 1 ( k[TlSi) } dt i?iV1+P2v2+P3v3+P4v4+P5vs $4 (4. lid) ^ 5 _r______^v,______jc[ri0t5gj] dt PiV1+P2v2+P3v3+P4v4+P5v5 €5 (4. lie) At each boundary the partitioning of the flux of silicon is given by [1 1 0 ,1 1 1 ] Xi/X2 - PiV1/P2V2 (4.12a) X2/X3 - P2V 2 /P3V 3 (4.12b) X3/X4 “ P3V 3 /P4V4 (4.12c) X4/Xs « P4V4/P5V5 (4.12d) Substitution of Eqs. (4.11) and (4.12) into Eqs. (4.9) and (4.10) results in five different equations to determine the global growth rate (K) for the total coating: 79 [i+ii+ii+ii+ii]2 k(ri3si) fA „ . _____ Si St Si Si______3 (4.13a) U W i V i 5 i W i . . k m ’" sl) (4.i3b , f ^2^1 . ^av a^3 ^av 4^< . ^av 5^5 i vtfi v;v2ea i4i*4*-*x*4**4*-i* *(w5/4si, (4.13c) if- t»3 *3 ^3 ^3______, y>tt. ... vs.*.. y»v»{», % % v , f . v , ? , I-!1 *-!1 *-!1*1*-!4 !* kiTiSi) .. {4 (1 (< (.______(4.13d) n«* ".s. v . £ « jy_ ‘r Sg ^ Sg r ^ Sg r Sg *______11’ * (rJ--»si> (4.13e) 13e, ^Ss V2is V,is v,(, 4.1.3 Growth Kinetics Figure 15 shows the fit to a linear regression plot for thickness versus the square-root of tine for all five of the titanium-silicide phases at 1423K. The error bars are one standard deviation, and no error bar is shown for points with a standard deviation smaller than the plotting symbol. eo The correlation coefficient was at least 0.98 for each silicide phase at every temperature, which indicates that the growth rates are parabolic. The apparent rate constant (kp') was determined for each phase by accounting for the incubation time (a) in the expression [113], Ci - a + (2kp 't)1/2 (4.14) The intrinsic rate constants were determined using Eqs. (4.4a-4.4e) and are reported in Table 9. The standard deviation for each intrinsic rate constant was estimated by assuming that the data fit a normalize t-distribution [142]. Figure 16 is an Arrhenius plot of the intrinsic rate constants for each phase, with the error bars representing one standard deviation. The activation energies for the diffusion-controlled growth of each layer are reported in Table 10. The interphase boundaries of the silicide phases in Fig. 13 are slightly non-planar. This is explained by the different orientations of the columnar grains. The diffusion constant is a tensor property, and the diffusion flux depends upon the orientation of the crystal, except for an isotropic solid [104]. TiSi2 and TiSi are orthorhombic, Ti5Si3 is hexagonal, and TisSi4 and Ti3si are tetragonal, so the diffusion rates depend on the orientation of each phase, which explains the slightly non-planar interphase boundaries. The orientations of the grains in each silicide layer were not determined, and the diffusion coefficients 81 Table 9. Intrinsic parabolic rate constants (kp [cm2/s]) measured by a titanium-silicon diffusion couple at 950, 1050 and 1150°C. 1150®C 1050°C 950°C TiSi, 1.44X10"10 4.03X10"11 1.42X10-11 TiSi 1.43X10"10 9.41X10”12 7.16X10**12 Ti5Si4 8.88X10“U 2.00X10"12 8.52X10"12 Ti5Si3 9.14X10-12 3.33X10"12 1.94X10-12 Ti3Si 3.40X10**12 1.72X10"12 3.07X10-13 OTiSf2 VTiSi nTi5Si4 OTi5Si3 AT13Si 601 — '— Ti3Si & Tl5SI3i I 1 l - 1 A I i ^ i6" 50 100 150 200 250 300 350 400 450 S quare Root Time (sec)^/^ Figure 15. A linear regression fit to the plot of the thickness of the TiSi2, TiSi, Ti5Si4, Ti5Si3, and Ti,Si layers versus the square root of time for the titanium- silicon diffusion couple 1150°C. The error bars are one standard deviation, and no error bar is shown for a standard deviation smaller than the plotting symbol. 82 Table 10. Activation energies for the growth of the five silicide layers in a titanium-silicon diffusion couple. TiSi, 167.000 J/mole TiSi 220.000 J/mole 168.000 J/mole 111.000 J/mole 176.000 J/mole U50C 1050C 950C I 3 t 1--- 1--- r 0.0006 0.0007 0.0008 0.0009 1/T (K'1) Figure 16. A linear regression fit to an Arrhenius plot’of log intrinsic parabolic rate constant versus inverse temperature for the TiSi?, TiSi, Ti5Si4, TisSi-j, and Ti3Si layers in the titanium-silicon diffusion couple. The error bars are the estimated standard deviation, with no error bar shown if the plot symbol is smaller than the standard deviation. 83 reported in this study represent averaged values for the prevailing grain orientations. Comparison of the measured intrinsic rate constants for Tisi2 growth in Table 9 with the values derived from the extrapolation of the diffusion data in Table 8 reveals that the measured values are an order of magnitude lower than the data reported by Revesz et al. [141], but only a factor of 2 or 3 lower than Hung et al. [140]. The measured activation energy (167,000 J/mole) for the growth of TiSi2 is approximately the same as that using the data of Hung et al. [140] (176,000 J/mole) but is significantly different from that using the data of Revesz et al. [141] (225,600 J/mole). Revesz et al. [141] measured the lateral growth rates of TiSi2 by evaporation of Ti onto Si02 islands on a preoxidized silicon wafer. Since TiO is more stable than Si02 [143], Sio2 would dissociate and the Ti film would absorb oxygen or form an oxide during the diffusion experiment. This may explain the difference between the values measured here and the data of Revesz et al. [141]. Hung et al. [140] measured growth rates at lower temperatures (475-550°: T/Tro » 0.464 to 0.422) than those used here, but the growth rates and activation energy for TiSi2 formation are about the same as the current study. In general, grain boundary diffusion is expected to be dominant at temperatures below half the melting point (in Kelvin) [10,104,105]. However, volume diffusion in the bulk would be expected to be dominant at the temperatures used in this study (T/Tm = 0.803 to 0.690). Volume diffusion in the bulk is dominant in certain other systems at normalized temperatures below that used in this study: NiO [10] (T/Tm = 0.653), Ag [104] (T/Tra * 0.70) and Cu3Si [43] (T/Tm = 0.656). Furthermore, Murarka and Fraser [145] used resistivity measurements to study the Tisi2 growth rate that results from the interdiffusion of titanium-silicon thin films at 500-800°C (T/Tm = 0.605 to 0.436), and observed an activation energy that is equivalent to the value determined in this study and by Hung et al. [140] at 475-550°C. The same mechanism for diffusion in TiSi2 is dominant over a * large range in temperature (T/Tm « 0.809 to 0.422), and is assumed to be bulk diffusion. The growth of compound layers in a diffusion couple by short-circuit paths results in the following morphology: (1) narrowly spaced, columnar grains, or (2) well-defined protrusions at grain boundaries that are in contact with the phase being consumed by the growing layer [105,146]. The cross-section of the diffusion couple in Fig. 13 revealed no protrusions in the TiSi2 grain boundaries at the TiSi/TiSi2 interface. Additionally, the columnar TiSi2 grains had widths in the range 10-40 /urn, which is at least an order of magnitude thicker than the grain width observed for Cu3Si via grain boundary transport [144]. For an average grain size of 20 /m and an effective boundary width of 3X10"8 cm 85 [104], the fraction of grain boundary area Is 1.5X10-5, an Insignificant value. The penetration volume ((Dt)1/2) is much larger than the grain size, which indicates that volume diffusion dominates grain boundary diffusion [105]. These results support the argument that volume diffusion is the controlling mechanism. However, the rates calculated by extrapolation of the low temperature data [140] (with short- circuit contributions) are twice the intrinsic rate constants determined in this study. The growth rates of TiSi in Table 8, that were calculated by extrapolation from the low temperature (T/Tra ** 0.297-.332) diffusion data of Chambers et al. [139], are much higher than the measured values. The measured activation energy (220,000 J/mole) is twice the value derived from Chambers et al. [113] (110,000 J/mole). For metals, the activation energy for volume diffusion is approximately twice the value of grain boundary diffusion [104]. Thus, the low temperature data of Chambers et al. [139] are most likely dominated by grain boundary transport, while the values measured in this study are for volume diffusion. Figure 13 shows that the columnar grains of the TiSi layer are an order of magnitude thinner than the columnar grains of the outer TiSi2 layer. Since the homologous temperature for TiSi (T/Tra = 0.772-0.664) is less than that for TiSi2 (T/Tm » 0.803-0.690), the observation of volume diffusion dominated growth for TiSi is consistent with the previous conclusion that the growth of Tisi2 is also controlled by volume diffusion. No diffusion data are available for Ti5Si4, Ti5Si3 and Ti3Si, and further comparison of these results was not possible, but their parabolic thickening rates are also assumed to be controlled by volume diffusion. Thus diffusion coefficients were calculated by substituting the intrinsic growth rate constants into Eqs. (4.6a-4.6e). Figure 17 shows an excellent linear regression fit to an Arrhenius plot for the calculated diffusivities for TiSi2, TiSi, Ti5Si4, Ti5Si3 and Ti3Si. Based on this fit to the data, the interdiffusion coefficients (D[cm2/s]), Q[KJ/mole)) for Tisi2, TiSi, Ti5Si4, Ti5si3 and Ti3Si are: D[TiSi2] - (1.25X10“4)exp(-169.6/RT) 4.15a) D[TiSi] - (3.25X10”3)exp(-197.2/RT) 4.15b) D[Ti5Si4] - (6.67X10”4)exp(-201.4/RT) 4.15c) D[Ti5Si3) - (2.05X10“6)exp(-160,7/RT) 4.15d) D[Ti3Si] » (4.89)exp(-297.2/RT) 4.15e) 4.2 Growth of Undopod Silicide coatings This study attempted to clarify the differences in layer growth kinetics between pack aluminizing and siliconizing, and establish baseline data for the growth of silicide coatings. To study the relationship between the stability of the halide activator and the growth rate, three activators with different stabilities were used with a pure Si source to form silicide coatings on pure Ti. The partial 87 pressures of the Si fluorides for each activator, which decide the magnitudes for the gas-phase fluxes, were determined by computer-assisted thermodynamic calculations. The growth rates of the coating for the three activators were experimentally determined, and compared with the rate of solid-state diffusion to clarify the rate-limiting step for siliconizing Ti by HAPC. 4*2.1 Thermodynamics of siliconising The normalized standard Gibbs energies of formation for the three activators presented in Table 11 [147] show large differences in magnitude. However, all three of these activators (unlike e.g. A1C13 or NH4F) maintain a condensed phase over the range of temperature used in this study. The Stepsol computer program [133] was used to determine the equilibrium partial pressures of the silicon fluoride vapor species by considering equilibrium between pure Si, the halide activator and an A1203 filler. Table 12 shows the typical output of such a calculation for a MgF2-activated pack at 1150°C. The high vapor pressure of aluminum fluoride species indicates that the A1203 filler is not truly inert. The implications of high Al fluoride vapor pressures are addressed later. Figures 18, 19, and 20 show the respective equilibrium partial pressures of the silicon fluoride species for MgF2- , A1F3- and CuF2-activated packs at 1000K to 1600K. Consistent with the thermodynamic data in Table 11, the 88 1150C 1050C 950C .*9 CO M .TiSi o6 i»10 TiSi. uc TLSi- •H Ti.Si. .*13 0.0006 0.0007 0.0008 0.0009 1/T (K*1) Figure 17. A linear regression fit to an Arrhenius plot of the calculated diffusion coefficient versus inverse temperature for TiSi2, TiSi, TicSi4, TisSi3, and Ti3Si from the diffusion couple measurements. Table 11. Standard Gibbs energies of formation normalized to one-half mole of molecular fluorine for MgF2, A1F3, and Cu F2 at 1150°C [143], i.e.: (l/y)X + l/2Fj « (l/y)XFy Ag ° [J/mole] 1423K 1323K 1223K MgF2 -435,730 -447,060 -455,720 A1F3 -380,300 -388,840 -397,420 CUF2 -175,426 -179,500 -183,400 89 Table 12. The calculated partial pressures for equilibrium of silicon, MqF2 activator and A120 3 filler at 1150°C. Partial Partial Pressure Pressure Ar 9.996X10"1 Si3 1.677X10"13 Mg 1.746X10 Si02 6.430X10”14 A1F 9.642X10"® F 1.799X10” 14 SiF4 4.283X10“ AlO 2.082X10"15 SiO 3.379X10“® Mgo 1.411X10-1® SiF3 2.717X10"® o 3.675X10”1® A1F 3 2.262X10“® aio2 2.849X10" MgF 2 1.545X10“® °2 7.445X10-2® SiF2 1.131X10“® *2 8.39BX10”29 A1F 2 9.038X10“® FO 1.252X10-30 MgF 4.482X10“® F02 8.731X10-43 AlOFo 8•849X10"7 f 2o 2.808X10”45 SiF 9.464X10“® °3 2.946X10"47 SiOF2 6.285X10"® Mg2F4 3.288X10“® Si 1.437X10 AlOF 9.470X10"10 Al 9.288X10“ A12F6 1.402X10 Mg2 2.945X10“ A120 1.590X10"11 Si2 1.090X10 Table 13. The sum of SiF3, SiF2 SiF and Si partial pressures for packs activated with MgF2, A1F3 and CuF2 calculated at 1150°C, 1050*C and 950°c. 1150°C 1050°C 950®C MgF2 3.849X10"® 3.847X10“® 3.108X10"7 AIF 3 2.627X10“3 3.203X10”4 2.736X10“ CuF2 5.000X10"2 2.305X10” 9.268X10“ 90 M Ma n3 s SIF, o Q. O* SIF 3 SIF, 1000 1100 1200 1300 1900 1M0 Temperature (K) Figure 18. Thermodynamic calculation of the partial pressures of the silicon fluorides for pure silicon, MgF2 activator and A1203 filler at 1000K to 1600K. QJO in3 n StF. £ 0L "3 ‘■2a Cl SIF cn 3 1000 1100 1300 1900 1900 Figure 19. Thermodynamic calculation of the partial pressure of the silicon fluorides species of pack consisting of pure silicon, A1F3 activator and A1203 filler at 1000K to 1600K. 91 M SF, £a «□ s SIF, o ••eo CL SIF 1000 1100 1200 1300 1400 1900 1100 Temperature (K) Figure 20. Thermodynamic calculation of the partial pressure of the silicon fluorides species of pack consisting of pure silicon, CuF, activator and A1203 filler at 1000K to 1600K. equilibrium vapor pressures of the Si fluorides decrease with increasing activator stability. In principle, the deposition of Si from the halide vapors onto the substrate may occur by reactions involving exchange (Eq. 1.3), dissociation (Eq. 1.4), disproportionation (Eq. 1.5) or an alkali vapor species (Eq. 1.6). The lower molecular weight of si compared to titanium predicts a weight loss for the exchange reaction, which was not observed. Since the dissociation reaction is not thermodynamically favorable, the disproportionation and alkali vapor reactions, Eqs. (1.5) and (1.6), are assumed to support the deposition of silicon from the gas phase. Figures 18 to 20 show that the ranking of silicon fluorides is SiF4, SiF3, SiF2 and SiF. Reactions (1.5) and (1.6) show that the deposition of Si can not result from SiF4, which is the Si fluoride in the highest oxidation state, so siF4 can only serve to return fluorine vapor back to the pack. Therefore, the flux of Si in the gas phase to the substrate results from the sum of the fluxes for SiF3, siF2, SiF and Si. The gas-phase flux for a given species is directly proportional its the partial pressure gradient. The partial pressures for the important Si fluorides for the three activators at 950°C, 1050°C and 1150°C are given in Table 13. 4.2*2 Characterisation of As-Coated Substrates The EDS profiles for silicide coatings grown using a MgF2- and AlF3-activated pack at 1150°C for 12 hours in Fig. 21 show a five-layer coating consisting of TiSi2, TiSi, Ti5Si4, Ti5Si3 and Ti3Si adjacent to the substrate. Thus, all of the intermediate equilibrium phases in the Ti-Si system are present [57]. The growth of these coatings is analogous to a Ti-Si diffusion couple, except that Si is provided by the gas phase in HAPC. Figures 22a, 22b, and 22c, respectively, show optical micrographs of the cross- sections of coatings grown using MgF2, A1F3 and CuF2 activators for 12 hours at 1150°C. The five-layered TiSi2/TiSi/Ti5Si4/Ti5Si3/Ti3Si structure of each coating is obvious upon etching, as indicated in Fig. 22. Solid-state diffusion is an atomistic process and inward diffusion results in vacancy condensation at the coating/gas 93 interface, which does not provide for a mechanism to form porosity in the coating [104,105]. Furthermore, cracking and grain pullout can result from polishing a brittle material, which gives the appearance of porosity [27]. Thus, any porosity or cracks observed in these coatings are an artifact of specimen preparation and etching. X-ray diffraction of the coating surface detected TiSi2 (JCPDS# 35-785, space group #70), and also minor A1F3 and a- A1203 peaks (Appendix A, Fig. 145a). Additionally, some minor MgF2 peaks were observed on the coatings grown using a MgF2 activator, due to condensation of the MgF2 salt during the coating treatment (Fig. 145b). Vigorous cleaning of as- coated substrates in boiling water did not change these results. However, CuF2 was not detected on the surface of coupons grown with CuF2-activated packs. The lower stability of CuF2 and low Cu vapor pressure in comparison to AIF3 may explain this observation. Examination of the surfaces of the coatings at high magnification by SEM revealed a porous, aluminum-rich layer. Guille et al. [52] also observed an adherent layer of "porous activator scale" after siliconizing Ti using an A1F3 activator. Additionally, activator-base deposits are observed at the surface of aluminide coatings grown by halide salts that are a condensed phase during the coating treatment [66,85]; i.e., the condensed activator mechanism in Figs. 3 and 4. Table 12 shows that the partial pressures of aluminum 94 100 00 TiSlj TiSi n5SI4 < •TT3Si M 00 E 70 5si3 o o — o —~ o 0 - 0 - 0 ■■ 0. S 00 \ c o 50 Silicon Profile b-o-o-o> o-o-c S 11 s. 40 B a s e E 30 cS 20 \ 10• (a) 0 to 20 30 40 SO 60 70 00 90 100 Distance From Outer Edge of Coating 0*m) 100 90 TISI, TiSi |n 5St4 K4n3si S ? 00 rn5si3 I 70 5> so o SO Silicon Profile Titanium '§ « CL S u b s tr a te B 30 5 JO 10 (b) 0 H h 0 10 20 30 40 SO 60 70 60 90 100 Distance From Outer Edge of Coating (jim ) Figure 21. Quantitative EDS profiles of silicon in the cross-section of the multilayered silicide coating after 12 hours at 1150°C: (a) 10 wt% Si, 2 wt% MgF2 and A1203 filler, (b) 10 wt% Si, 2 wt% A1F3 and A1203 filler. 95 Figure 22. Optical micrograph of the cross-section of a silicide coating grown for 12 hours at 1150°c. The Kroll etch was used and the five TiSi2/TiSi/Ti5Si4/Ti5Si3/Ti3Si layers are marked on the micrograph. (a) 10 wt% silicon, 2 wt% MgF2 and A1203 pack, (b) 10wt% silicon, 2 wt% A1F3 and A1203 pack, (c) 10wt% silicon, 2 wt% CuF2 and A1203 pack. fluorides are quite high for a MgF2-activated pack, which is the most stable of the three activators. During the growth of the coating, the condensation of aluminum fluoride on the substrate is expected thermodynamically, and this condensation could entrap some inert A1203 at the surface, as marked in Fig. 22. The most intense XRD peaks for A1F3 and A1203 were observed on coatings formed by the least stable C u F2 activator, as expected. The EDS profiles of the cross-sections of the these coatings did not detect any aluminum within the coating layers. Since diffusion in Tisi2 and Tisi is dominated by migration of the smaller silicon atom [137,139], these coatings are grown inward. Therefore, the condensed activator skin would be expected to remain at the coating surface, as an inert marker. After sanding the coating grown from an AlF3-activated pack, XRD detected the crystal structure of the remaining phases: Tisi (JCPDS# 17-424, S.G. #62), Ti5Si4 (JCPDS# 27-907 S.G. #92) and Ti5Si3 (JCPDS# 29-1362 S.G. #193), with Ti3Si only seen by SEM (Appendix A, Fig. 146). 4.2.3 Growth of silicide Diffusion coatings Figures 23, 24 and 25 show the respective parabolic thickening plots for MgF2-, A1F3- and c u f 2-activated packs at 1150°C. The correlation coefficient for almost every plot was 0.95 or greater, which indicates a parabolic growth rate. Apparent parabolic rate constants were determined for each layer at each of the three temperatures studied using 97 Table 14. Intrinsic parabolic rate constants (k- [cm2/s]) determined from the experimental data for a MgF2- activated pack at 950, 1050 and 1150°C. 1150°C 1050°C 950®C Tisi, 1.19X10”10 8.85X10"11 4.96X10"11 TiSi 1.13X10“10 2.83X10 8.27X10”12 TisSi4 6.92X10”11 8.87X10 1.76X10"12 Ti5Si3 7.66X10-12 5.61X10"12 9.86X10”13 Ti3Si 5.11X10“12 2.23X10 1.46X10“ OTiSU VTiSi DTicSiA OTieSi-* ATi-tSi SO t - 4 5 " 4 0 - 3 5 - TiSi 3 0 - n vn 2 5 - TiSi 5 o 20 - 1 5 - 1 0 - 5 - 0 50100 150 200 250 300 350 Square Root Time (Sec) Figure 23. A linear regression fit to the plot of the thickness of the TiSi2, Tisi, Ti5Si4, Ti5Si3, and Ti3Si layers versus the square-root of time in a coating grown at 1150°C from a 10wt% silicon, 2 wt% MgF2 and A1203 filler pack. 98 Table 15. Intrinsic parabolic rate constants (kp [cm2/s]) calculated from the experimental data of an A1F3- activated pack at 950, 1050 and 1150°C. 1150°C 1050°C 950°C TiSi, 7.10X10"11 4.17X10“n 7.49X10“12 TiSi 7.91X10*11 3.48X10"11 4 • 28X10 Ti5Si4 2.27X10*11 1.26X10"11 1.74X10"12 TisSi3 1.02X10"11 4.51X10-12 1,78X10 Ti3Si 6.10X10”12 2.21X10"12 7.10X10“ OTiSio VTiSi DTI 40 30-- TlSi 01 m at c TiSi o 10-- 5-- 0 50 100 150 200 250 300 350 Square Root Time (Sec) Figure 24. Plot of the thickness of the TiSi2, TiSi, Ti5Si4, Ti5Si3, and Ti3Si layers versus the square-root of time for a coating grown at 1150°C with a pack composed of 10wt% silicon, 2 wt% A1F3 and A1203 filler. 99 Table 16. Intrinsic parabolic rate constants (k. [cm2/s]) calculated from the experimental data for a CuF2- activated pack at 950, 1050 and 1150°C. 1150°C 1050°C 950°C TiSi, 6.02X10"11 -11 5.64X10"12 TiSi 7.23X10"11 -11 4.03X10"12 Ti5Si4 3.14X10“U -11 2.48X10”12 TisSi3 5.55X10-12 -12 1.01X10"12 TijSi 6.25X10"12 “12 5.33X10-13 OTlSi2 VTlSi □ TlgSi^ OTi5Si3 ATi3 Si 25 y 20-* 1 15-* n a)« c u 10-* £ 5* oil 0 50 100 150 200 250 Square Root Time (sec)^/^ Figure 25. Plot of the thickness of the TiSi2, Tisi, Ti5Si4, Ti5Si3, and Ti3Si layers versus the square-root of time for a coating grown at 1150°C with a pack composed of 10wt% silicon, 2 wt% cuF2 and A1203 filler. 100 Table 17. Activation energies (Q [J/mole]) for the growth of each layer of the multi-layered coating for MgF2-, AIF3- and CuF2-activated packs. MgF2 aif3 cuf2 TiSi2 66,700 165,000 172,000 TiSi 189,000 213,000 212,000 TisSi4 264,000 188,000 185,000 TisSi3 151,000 126,000 122,000 Ti3Si 89,700 156,000 179,000 1150C 1050C 950C TiSi, S 0 , O. _r c 13 TLSi ■TiSi 2 o U <0 cd OS o ,-12 TLSi, o x» TLSi ol. ,-13 0.0006 0.0007 0.0008 0.0009 l/T (K'1) Figure 26. A linear regression fit to an Arrhenius plot of log intrinsic parabolic rate constant versus inverse temperature for the TiSi2, Tisi, Ti5Sia, Ti5Si3, and Ti3Si layers in a coating grown by a MgF2-activated pack. 101 1150C 1050C 95 OC -10 TiSi, Ti.Si 2Jea a; u ,-12 o JO TLSi Cu ,* 1 3 0.0006 0.0007 0.0008 0.0009 l/T (K-1) Figure 27. A linear regression fit of an Arrhenius plot of log intrinsic parabolic rate constant versus inverse temperature for the TiSi2# Tisi, Ti5Si4, Ti5Si3, and Ti3Si layers of an A1F3- activated pack. 102 1150C 1050C 950C "e p , ,-10 :isi eac TiSi, Ti,sr u ,-12 3 Ti,Si. O. TLSi ,-13 0.0006 0.0007 0.0008 0.0009 1/T (K*1) Figure 28. A linear regression fit to the Arrhenius plot of log intrinsic parabolic rate constant versus inverse temperature for the TiSi2, Tisi, Ti5Si4, Ti5Si3, and Ti3Si layers in a coating formed by of a CuF2-activated pack. 103 Eq. (4.14). The Intrinsic rate constants were determined using Eqs. (4.4a-4.4e). Tables 14, 15 and 16 report the intrinsic parabolic rate constants for each layer of the MgF2-, A1F3- and CuF2-activated packs, respectively. Figures 26, 27 and 2B are, respectively, the Arrhenius plots of the intrinsic parabolic rate constants for the five layers in each coating for the MgF2-, A1F3- and CuF2- activated packs. The activation energies for the growth of each layer of the coatings are given in Table 17. The intrinsic rate constants for the growth of each of the five silicide layers using A1F3 and CuF2 activators are less than a factor of two lower than the solid-state diffusion rate in Table 9. Then some slowing by the gas- phase diffusion step results in a slightly lower growth rates for the diffusion coatings. Less than 15% difference was observed between the activation energy for the growth of each phase in the five-layered coating by A1F3- or CuF2- activated packs in Table 17 and for the Ti-Si diffusion couple in Table 10. Since the values of the activation energy and intrinsic rate constant are almost the same as the results obtained from the diffusion couple, solid-state diffusion dominates the growth of silicide diffusion coatings formed from the A1F3- and CuF2-activated packs. Furthermore, Figs. 151a and 151b in Appendix B show that the thickness ratios for the layers become a constant at steady state, as expected for the growth of multiple layers by 104 solid-state diffusion [109-112]. The incubation tine (Eq. 4.14) produced the deviations from a constant thickness ratio that were observed in the initial period of coating growth. The activation energy for TiSi2 grown by the MgF2- activated pack is lower than the value determined from the diffusion couple. The outer surface of Tisi2 observed in Fig. 22a is non-planar, which nay result from gas-phase transport as the rate-limiting step [148]. Bianco and Rapp [77] demonstrated that when a porous, refractory shield was used to prevent pack entrapment during the aluminizing of nickel-based alloys, the activation energy for growth of the coatings was lowered. The porous refractory inhibited the gas-phase transport, and changed the rate-limiting step from solid-state diffusion control to partial control by gas- phase diffusion. Levine and Caves [66] also reported that the use of a more stable activator resulted in lower vapor pressures for the A1 halides and an activation energy less than the solid-state diffusion value. These kinetics were controlled by gas-phase diffusion of the volatile halides to the substrate [66]. Since the activation energy for TiSi2 growth is much lower than the solid-state value, the kinetics are probably partially controlled by gas-phase diffusion of the si fluorides. The vapor pressures of the fluorides, as summarized in Table 13, are much lower for the MgF2 activator than for the A1F3- and CuF2-activated packs, which may result in partial control by gas-phase diffusion. However, the growth rate of all five silicide phases grown by a MgF2-activated pack are approximately equal to the diffusion couple growth rates at 1150°C, which indicates that solid-state diffusion may be rate-controlling at 1150®c. The MgF2-activated growth rates for all phases are significantly different than the rates of the diffusion couple at 1050 and 950°C, which indicates that gas-phase diffusion is limiting. The large difference between the activation energies measured from a MgF2-activated pack and the diffusion couple for the Tisi, Ti5Si4, Ti5Si3 and Ti3Si phases shows that the inner layers of a multi-layer scale are also affected by the limiting of gas-phase diffusion. According to Table 13, the partial pressures of Si fluorides in the gas phase must be at least 10~S atm. for solid-state diffusion to control the growth kinetics in this system. The thickness ratios for silicide coatings grown by a MgF2-activated pack in Fig. 151c of Appendix B do not closely approach a constant value at steady state, which should be observed for diffusion-controlled growth [109- 112]. The thickness ratios are closest to a constant value at 1150°C. These results indicate that the growth is affected by gas-phase diffusion, but solid-state diffusion may be rate-limiting at 1150°C, in agreement with the previous discussion. However, the Arrhenius plot in Fig. 26 shows that the data falls on a straight line, which should 106 not be observed for a complete change In growth mechanism. Growth kinetics should be measured at temperatures above 1150°C to resolve a change from solid-state to gas-phase diffusion. Comparison of Tables 14, 15 and 16 shows that the intrinsic rate constants for the growth of TiSi2 are higher when a less stable MgF2 activator is used to form the coating. In fact, the rate of MgF2-activated TiSi2 growth exceeds the diffusion couple rate at 1050 and 950°C. The high vapor pressures resulting from the use of a less stable activator may result in some venting of vapors from the retort, which could explain the slower kinetics. Guille et al. [52] varied the amount of A1F3 activator UBed to siliconize CP titanium, and observed that for contents of 2 wt.% or larger, the growth rate did not change. Therefore, the escape of gases from the sealed retort was assumed to have a negligible effect on the coating kinetics. The lesser amount of A1F3 and A1203 condensation resulting from the use of the more stable MgF2 activator may also contribute to the slower growth kinetics. However, the observation of larger TiSi2 growth rates for a MgF2- activated pack may result from experimental uncertainty in measuring the thickness of the non-planar TiSi2 phase. Furthermore, if the Si fluoride vapor pressures influence the interdiffusion constant for inward si diffusion in Tisi2 through the vacancy concentration, then a faster TiSi2 107 growth rate would be expected for a lower partial pressure of Si fluoride, in agreement with the results [10,104]. 4.2.4 Global Growth hates The global growth rates for the entire diffusion coating and diffusion couple were determined from the intrinsic rate constants of each layer by Eqs. (4.13a- 4.13e). The calculated global rate constant determined from each layer vary by less than a factor of two at a specific temperature for each of the diffusion coatings in Table 18 and diffusion couple in Table 19. However, experimental error associated with measuring the thin Ti5si3 and Ti3si layers results in a larger variation for the calculated global rate constant. The global rate constants measured for the overall coating, given in Table 20, agree well with global rate constants calculated from the intrinsic rate constant of the individual silicide layers. The self- consistency between the measured and calculated global rate constants for the five-layered diffusion coating and the diffusion couple indicates a valid application of the theoretical treatment to the experimental results. 4.3 Growth of Titaniua-Borids Diffusion Coatings The growth rates of pure Ti-boride coatings weredetermined to study the influence of boron additions on the growth kinetics for the B-doped silicide coatings. The kinetics were measured at 1150°C, 1050°C and 950°c using a high stability HgF2 activator and the less stable A1F3 10S Table 18. Global growth rate (cm2/s) for the overall coating that are calculated from the intrinsic rate constant for MgF2-, A1F3- and CuF2-activated packs at 950°C, 1050°C and 1150°C. MgF2 1150°C 1050°C 950°C TiSi2 2.63X10“ 1.65X10"10 8.26X10-11 TiSi 5.54X10-10 1.47X10“10 5.33X10-11 Ti5Si4 6.34X10"10 1.31X10-10 3.72X10"11 Ti5SiJ 4.45X10"10 3.96X10"10 7.55X10"11 Ti3Si 2.24X10”10 1.38X10-10 7.27X10-11 AlFi 1150°C 1050°C 950#C TiSi, 1.81X10“ l.OOXlO-10 1.77X10* TiSi 3.31X10-10 1.35X10-10 1.92X10* TisSi4 2.06X10"10 1.49X10"10 1.76X10* Ti5Si3 4.32X10-10 2.47X10-10 8.05X10* Ti3Si 1.98X10-10 8.08X10"11 2.34X10* CuF5 1150°C 1050°C 950WC TiSi, 2.02X10-10 8.49X10-11 1.46X10-11 Tisi 2.73X10*10 1.36X10“10 1.73X10”11 Ti5Si4 2.02X10"10 1.05X10-11 2.39X10"11 Ti5Si3 1.78X10-10 7.28X10-11 4.11X10"11 Ti3Si 1.60X10"10 1.04X10”10 1.33X10-11 109 Table 19. Global rate constants (cm2/s) for the overall growth of the titanium-silicon diffusion couple calculated from the intrinsic rate constants at 950°C, 1050°C and 1150’C. 1150°C 1050*0 950*C Tisi, 3.55X10"10 8.28X10 2.18X10"11 TiSi 6.23X10"10 3.93X10 4.47X10"11 Ti5Si4 7.50X10"10 3.12X10 2.62X10"11 TicSig 4.04X10-10 2.84X10 3.89X10-11 TijSi 1.42X10"10 9.60X10 2.90X10"11 0 E 2 s BF £ 3 (0 BF V) -4 9> (0 -6 BF •co ■ e Q_ O) O 8 -10 1000 1100 1200 1300 1400 1500 1600 Temperature [K] Figure 29. Thermodynamic calculation of the equilibrium boron fluoride pressures for a pack comprising of B, MgF2 and A1203 at 1000 to 1600K. 110 Table 20. Global rate constants (cm2/s) for the overall growth rate of the five-layered silicide coatings measured at 950°C, 1050° and 1150*C for MgF2, A1F3 and Cu F2 activators, and for the titanium-silicon diffusion couple. Activator 1150°C 1050®C 950°C MgF2 3.82X10”*° 1.44X10"10 8.89X10**11 A1F3 2.33X10“10 2.01X10-10 2.04X10”11 Cu F2 2.00X10-10 1.21X10“10 1.78X10-11 Ti-Si Diffusion 4.06X10“*® 1.78X10"10 3.29X10"11 Couple E T5. BF, *£ 3 CO CO 0) ^ or 1000 1100 1200 1300 1400 1500 1600 Temperature [K] Figure 30. Thermodynamic calculation of the equilibrium boron fluoride pressures for a pack comprising B f A1F3 and A1203 at 1000 to 1600K. Ill activator (Table 3). 4.3.1 Thernodynenics of Boriding Figures 29 and 30 show the partial pressures for the B fluorides that are calculated using the STEPSOL computer program for packs comprised of either 4 wt.% B, 2% MgF2, and A1203, or else, 4% B, 2% A1F3, and A1203. The vapor pressures of the B fluorides are much higher for the less stable A1F3 activator, in agreement with the previous results for siliconizing. Similarly, the partial pressures for the Al- and Mg-fluorides were high enough to expect salt deposits by the condensation mechanism (Fig. 3) [66,85]. Since a weight loss was not observed for the deposition of B, the exchange mechanism is not important (Eq. 1.3), and the dissociation reaction (Eq. 1.4) is not thermodynamically favorable. Therefore, the deposition of boron from the gas phase is assumed to occur by the disproportionation (Eq. 1.5) and alkali vapor reactions (Eq. 1.6). Figures 29 and 30 show that the ranking of boron fluorides is BF3, BF2, BF and B2F4. The deposition of boron does not result from BF3 for the reactions described by Eqs. (1.5) and (1.6), and BF3 is the return species. The gas-phase flux for deposition is the sum of partial pressures for B, BF, BF2 and B2F4, which are given in Table 21. 4.3.2 Kinetics of boriding Figures 31 and 32 are optical micrographs of boride coatings grown by MgF2- and AlF3-activated packs, 112 Table 21. The sum of BF2, BF, B and B2F4 partial pressures calculated for packs activated by MgF2 and A1F3. 1150°C 1050°C 950°C AIF3 5.62X10"4 6.08X10"® 4.53X10"6 MgF2 2.20X10"5 2.40X10 1.79X10"7 Figure 31. Optical micrograph of the cross-section of a boride coating grown by a pack comprised of 4 wt.% B, 2% MgF2, and A1203 at 1150°c for 12 hours. 113 50 pm Figure 32. Optical micrograph for a boride coating grown by a pack comprised of 4 wt.% B, 2% A1F3, and A1203 at 1150°C for 12 hours. The TiB2 and TiB phases are marked in the cross-section. respectively. A two-layer coating was produced by each pack, with an outer TiB2 and an inner TiB layer. The thickness of the inner TiB layer was very irregular, with thin regions and long, rod-like protrusions. Singhal [61] observed the same morphology for Ti-boride coatings grown on CP-titanium and a Ti-6A1-4V alloy by the HAPC method. Figures 33 and 34 show a linear regression fit for a plot of the TiB2 and TiB layer thicknesses for MgF2- and AlF3-activated packs, respectively. The error bars are one standard deviation, with large error bars shown for the nonuniform TiB layer. The correlation coefficient for the TiB2 layer was usually 0.95 or better, which indicates parabolic kinetics. Despite the large variation in thickness observed for the inner TiB layer, a parabolic fit 114 was assumed for this layer. The apparent rate constants were determined by correcting for the incubation time (Eq. (4.14)). Equation (2.25) was used to determine the intrinsic parabolic rate constants from the measured apparent rate constants, which are reported in Table 22. Samsonov and Latysheve [149] determined a grain boundary diffusion coefficient for TiB2 by measuring the recrystallization kinetics. No diffusion data were found in literature for TiB. A theoretical calculation for the growth kinetics of TiB2 was made using Eq. (4.5) at. 1150°C, 1050°C and 950°C, kp[TiB2] = D{-lnaB[TiB2/TiB]> (4.16) where D is the grain boundary diffusivity for TiB2, and aQ is the activity of boron at the TiB2/TiB interface. The results of this calculation are reported in Table 23. Figure 35 is an Arrhenius plot for the parabolic rate constants for the TiB2 and TiB layers grown by either a MgF2- or AlF3-activated pack. The activation energies for the growth of the boride layers are reported in Table 24, with the activation energy determined for TiB2 growth from the data of Samsonov and Latysheve [149]. The differences between the intrinsic rate constants for the TiB layers grown by the A1F3- and MgF2-activated packs are less than a factor of 2 or 3. However, large differences in activation energy were measured for the TiB layers. Since wide variations in thickness were observed for the TiB layers, an 115 Table 22. Intrinsic parabolic rate constants (kp [cm2/s]) determined for the growth of the dual-layer TiB2/TiB coating by packs activated with either MgF2 or A1F3 at 950°C, 1050#C and 1150«C. 1150°C 1050°C 950°C AlFi TiB, 3.08X10"12 2.49X10”12 4.52X10-13 TiB 1.22X10-11 8.28X10"12 1.58X10'13 MgF2 TiB, 1.25X10**11 9.50X10"12 1.28X10-12 TiB 1.87X10"11 1.37X10-11 5.69X10"13 16 *: TIB-11S0C 12*i J3 < — TlB*95dC ()T 50 100 150 200 250 300 350 Square Root Time [sec1/2] Figure 33. A linear regression fit for the plot of the average TiB2 and TiB thickness versus the square root of time for a MgF2-activated pack at 950°C, 1050°C, and 1150°C. The error bars represent one standard deviation. 116 Table 23. Intrinsic parabolic rate constants (k_ [cm2/s]) for TiB2 growth at 950“C, 1050°C, andFH 5 0 ttc that are calculated from the grain boundary diffusion data reported by Samsonov and Latysheve [149]. kp [cm2/s] 1150°C 9.50X10'11 1050°C 6.73X10*11 9S0“C 4.49X10-11 TiB-1150C E ,ti. TIB-1050C II---- — 0 50 100 150 200 250 Square Root Time [sec1/2] Figure 34. A linear regression fit for the plot of the TiB2 and TiB layer thickness versus the square root of time for an AlF-,-activated pack at 950°C, 1050°C, and 1150°C. The error bars represent one standard deviation. 117 Table 24. Activation energies (Q [J/mol]) determined from the intrinsic parabolic rate constants for the dual-layered growth of TiB2 and TiB by an AlF3- and MgFj-activated pack. Activation energy determined from the intrinsic rate constant for TiB2, which was calculated from the grain boundary diffusion data of Samsonov and Latysheve [149]. A1F3 MgF2 Literature Data [149] TiB2 141,000 168,000 54,200 TiB 320,000 54,200 1150C 1050C 9S0C -9 1 0 ~ i f ~ N E TIB, MgfyActivated u ...... c (0 4 -» ■11 V) 1 0 g . TIB^, MgFz-Actlvated c o (J Figure 35. A linear regression fit for an Arrhenius plot of the intrinsic parabolic rate constants versus inverse temperature for the TiB2 and TiB layers formed by either a MgF2- or AlF3-activated pack. 118 accurate measurement for the parabolic rate constant is difficult. In fact, the nonuniform thickness for the TiB layer indicates that the kinetics are not parabolic, which may explain the differences in the TiB growth kinetics for the AIF3- and MgF2-activated packs. The parabolic rate constants that were calculated from the grain boundary diffusion data of TiB2 (Table 23) are 4 to 8 times larger than the growth of TiB2 by a MgF2- activated pack, and 10 to 15 times larger for the A1F3- activated pack. The activation energy determined for the growth of TiB2 from the data of Samsonov and Latysheve [149] is less than half the values measured for the coatings grown by AIF3- and MgF2-activated packs. Therefore, the growth of the TiB2 diffusion coating layer does not occur by solid- state diffusion along the grain boundaries. The similarity in activation energy for the growth of TiB2 by an A1F3- and MgF2-activated pack indicates that the growth mechanism is the same. Since the intrinsic parabolic rate constants are lower than those determined for grain boundary diffusion, and since the measured activation energy is at least double that for grain boundary diffusion, the rate-controlling mechanism for growth of the TiB2 coating layer is bulk solid-state diffusion [104,105]. A change from solid-state diffusion control to gas- phase controlled kinetics was not observed for the growth of the TiB2 layer by the more stable MgF2 activator. The rate- 119 limiting step for the growth of silicide coatings by a MgF2- activated pack was gas-phase diffusion, and solid-state diffusion for the A1F3-activated pack. The gas-phase interdiffusion coefficient for the fluoride species {called i) in argon is approximated by [25], »l,»r " D » r < > W « l > 1/2 <«•«> where M is the molecular weight. Since the gas-phase diffusion coefficient for the lower molecular weight B- fluorides is higher than for the Si-fluorides, the gas-phase flux for boron was larger according to Eq. (2.2). The gas- phase flux of boron supplied to the surface of the workpiece was sufficient to maintain solid-state controlled kinetics for the lower B-fluoride vapor pressures. However, the growth rates for TiB2 by a MgF2-activated pack are 2-4 times higher than those for an AlF3-activated pack. The only obvious explanation for this difference is the non-parabolic growth kinetics for the inner TiB layer. Since the TiB growth kinetics affect the TiB2 layer, a model that is different from the one developed by Yurek et al. [110] may be required to describe the growth kinetics. 4.4 Boron-Doped Silioide Coatings One goal of this work was to produce silicide coatings with boron additions in a single reaction/processing step by the HAPC method. This task was accomplished by using thermodynamic calculations to direct the experiments. Preliminary work showed that a chloride-base activator 120 always produced a porous Ti silicide coating. The Tl chlorides are more stable than the Si chlorides, so that an exchange reaction (Eg. 1.3) produces a porous morphology. Since the Si fluorides are more stable than the Ti fluorides so that dense silicide coatings were produced (Fig. 22), fluoride-base activators were used for this work. The purpose of the boron additions is to dissolve into and improve the properties of the Si02 scale that forms on TiSi2 during high-temperature oxidation. B203-Si02 mixtures posses a higher CTE than pure Si02, which minimizes spalling during thermal cycling [39]. A more fluid B203-Si02 glass may heal cracks in the silicide coating during high- temperature oxidation [37]. Furthermore, boron forms a fast growing transient oxide that seals cracks, and is overgrown by Si02. 4*4*1 Development of B-Doped Silioide coatings on CP-Titanium Figure 36 shows the Si and B fluoride vapor pressures that were calculated for a pack comprised of pure si, pure B and a MgF2 activator. The vapor pressures for the B fluorides are a little higher than the si fluorides, but they are similar in magnitude, which indicates that the codeposition of si and B is possible. However, the coating produced by a pack comprised of 5 wt.% pure Si, 5% pure B, 2% MgF2 activator and A1203 filler was TiB2 with 2-3 at.% Si. These results were not changed by varying the filler, activator, or Si to B ratio of the pack. The growth of the boride coatings on CP-titanium (section 4.3) demonstrated that: (1) the low molecular weight for the B fluorides resulted in a high gas-phase diffusion coefficient, and (2) the gas-phase diffusion coefficient for the B fluorides is faster than the Si fluorides (Eq. 4.17). According to Eq. (2.2), a higher gas-phase diffusion coefficient results in a higher gas-phase flux for an equivalent gradient in partial pressure. The gas-phase fluxes for the B fluorides dominated those for the Si fluorides and resulted in the formation of a boride coating. TiB2 has poor oxidation resistance at high temperature [41], and is not the desired result. Codeposition has been achieved by using a binary alloy powder as the masteralloy, which lowers the thermodynamic activity for the dominant element and produces comparable gas-phase fluxes for the halides [35]. Figure 37 is a plot of the activity for the Si-B system at 1400K from the thermodynamic data of Dirkx and Spear [150]. The decrease in the B activity is less than 1 order in magnitude, with the exception of a very dilute Si-B alloy. Figure 38 is a plot of the Si- and B-fluoride vapor pressures for a pack comprised of pure Si, a SiB3 compound and MgF2 activator. The SiF4 and SiF3 vapor pressures are slightly larger than BF and BF2, which improves the possibility for codeposition. However, a homogeneous boride coating was the only result 122 SiB ' 3 > OjOOO 0.100 0.200 0300 0.400 0.900 a *00 0.700 0.000 0.(00 1000 Atomic Percent Silicon Figure 36. Partial pressures of Si and B fluorides calculated for a pack comprised of pure Si, pure B, and a MgF2 activator at 1000-1600 K. J M o. 3 to on L. M Q. "o •o ■ e 0. cn 3 SIF, sir 1000 1100 1200 1300 1400 1900 1000 Temperature (K) Figure 37. The Si and B activity in the Si-B system calculated at 1400K from the data of Dirkx and Spear [150]. A is a single-phase SiB,2 field, B is a two-phase SiB12-siB6 field, and C is a two-phase SiB6-SiB3 field. 1000 MOO 1200 1300 1400 1000 1000 Temperature (K) Figure 36. Partial pressures for the Si and B fluorides calculated for a pack comprised of pure Si, a SiB3 compound, and a HgF2 activator at 1000-1600 K. for a pack comprised of pure Si and SiB3 source. The partial pressures for the Si-fluorides are higher, but the higher gas-phase diffusion coefficients for the B-fluorides (Eg. 4.17} result in higher gas-phase fluxes for the B- fluorides by Eq. (2.2). The codeposition for Si and B cannot be predicted by designing a pack chemistry that produces equivalent vapor pressures. Therefore, the B activity must be lowered beyond the limits of the Si-B system to achieve codeposition. These observations indicate that a more stable boride compound must be used to lower the chemical activity of boron beyond the limits of the Si-B system. Since the gas- phase fluxes for the B fluorides dominate when the vapor pressures for the Si fluorides are higher than the B 124 fluorides, a large decrease in the partial pressures for the B fluorides with respect to the Si fluorides is required to achieve codeposition. Figure 39 shows that the activity for boron in the Ti-B system is at least 4 orders in magnitude lower than the Si-B system. The calculated vapor pressures for the B-fluorides are much lower than the Si-fluorides for a pack comprised of pure Si, a TiB2 compound and MgF2 activator (Fig. 40). Figure 41 demonstrates that success was achieved using a pack composed of 7 wt.% pure Si, 6t TiB2, 2% MgF2 activator and A1202 filler at 1150°C for 12 hours. Note the excellent adherence of the coating at the sharp corner in Fig. 41a. The microprobe trace in Fig. 41c indicates that the coating is a silicide, with all of the boron localized at the surface in a 1-2 pm layer. TiSi2 has no solubility for boron [152], and boron was not detected in the bulk of the silicide coating (Fig 41a), as expected. The surfaces phases detected by XRD (Appendix A, Fig. 147a) were TiSi2 (JCPDS# 35-785, space group #70) with minor peaks of TiB2 (JCPDS# 35-741), which verifies the microprobe results. Trace amounts of condensed activator salt were detected at the surface of the coating by XRD and SEM/EDS. The SEM micrograph of the as-coated surface in Fig. 42 shows that the TiB2 is present as discontinuous patches. SEM/EDS verified that the protrusions in Fig, 42 are TiB2. 125 Ti+TiB TiB TiB^ + TtB, B *1 0 < 01 O ■1 Ti -1 -1 + + + -i— h h— !— 4- .000 0.100 0.200 0J00 0.400 0.900 0.900 0.700 0400 0.900 1.000 Atomic Percent Boron Figure 39. The Si and B activity in the Ti-B system calculated at 1400K [151]. 0.0 E 2 .0 >>_o ' 4) m -4.0 • ■ Figure 40. Partial pressures for the Si and B fluorides calculated for a pack comprised of pure Si, a TiB2 compound, and a MgF2 activator at 1000-1600 K. 100 SO- T3SI2 BO. TiSI n s si+ s E o 70- Silicon Profile 3 , 60- ,0 -0 — o ------o ------c "N o 30- s3 [ ^ ^ 40- no E 30- «3 20- 10- ■7 <- Boron Profile 0- ft—a i a l a — H-- a - 1-- a - 1 -----a-t— —• a ■ 0 10 20 30 40 30 00 70 BO (c) Distance From Outer Edge of Coating Gun) Figure 41. The polished cross-section of the B-doped silicide coating grown on CP-titanium by a pack composed of 7 wt.% Si, 6% TiBj, 2% MgF2, and A1203 at 1150#C for 12 hours, (a) optical micrograph, (b) optical micrograph of etched cross-section and (c) microprobe profile of Si and B. 127 Closer examination of the cross-section in Fig. 41a by • SEM/EDS detected thin layers of TisSi3 and Ti3Si adjacent to the substrate, in addition to TiSi2, TiSi and Ti5Si4 (Fig. 41b and 41c). Similar to the undoped silicide coatings shown in Fig. 22, the B-doped silicide coating is a five- layered silicide coating. The morphology for the B-doped silicide coating (Fig. 41a) is an improvement over the undoped silicide coating grown by the MgF2 activator (Fig. 22a). Since boron has negligible solubility in TiSi2, the diffusion of B into the silicide layer was not observed, and the B additions remained at the surface such as inert markers. Titanium diboride is more stable than TiSi2 [138,151], and the small quantity of boron deposited at the surface reduces TiSi2 to produce TiB2, 2B(s) + TiSi2(s) = TiB2(s) + 2Si(s) (4.18) The displacement reaction described by Eq. (4.18) creates free Si that diffuses inward to grow Ti-silicide, which may have improved the morphology of the B-doped silicide coating formed by a MgF2-activated pack. Further investigation is required to clarify the exact mechanism that improves the morphology for the B-doped silicide coatings. Titanium diboride is one of the most stable boride compounds, and has a low chemical activity for B. An increase in the chemical activity for boron in the pack would raise the B concentration of the coating. Table 25 compares the Gibbs energy of formation and boron activity 128 \ H 2 J < \ V I « 1 . ' Figure 42. A SEM micrograph of the surface of a B-doped silicide coating grown on CP-titanium by a pack consisting of 7 wt.% Si, 6% TiB2, 2% MgF2 and A1203 at 1150°C for 12 hours. E a s 3 0.2 ST "5 ‘■eo ^ BF CL C7I 3 1000 1100 1100 1«0Q two Temperature (K) Figure 43. Partial pressures of the Si and B fluorides calculated for a pack composed of pure Si, a CrB2 compound, and a MgF2 activator at 1000-1600 K. 129 Table 25. The activity of boron and Gibbs energy of formation (AG*[J/mol]) normalized to one mole of boron for SiB,, CrB2, TaBa and TiB2 at 1423K in equilibrium with the pure metal [150,151]. AG°[J/mol] Boron Activity SiB3 -1,180 0.656 CrB2 -7,730 3.10X10"2 TaB2 -23,470 2.51X10"4 TiB2 -35,840 3.20X10“® E o £ 3 « «n □ ’-5 o Q_ cn o 1000 1100 1200 1300 1400 1900 ieoo Tem perature (K) Figure 44. Partial pressures for the si and B fluorides calculated for a pack comprised of pure Si, a TaB2 compound, and a MgF2 activator at 1000-1600 K. 130 0 0 8 n ® l J Figure 45. Optical micrograph of the polished cross- section of a B-doped silicide coating on CP-titanium grown by a pack comprised of 7 wt.% Si, 6% CrB2, 2% MgF2 and A1203 at 1150°C for 12 hours. BOB £3S1 Figure 46. Optical micrograph of the polished cross- section of a B-doped silicide coating on CP-titanium grown by a pack comprised of 7 wt.% Si, 6% TaB2, 2% MgF2 and A1203 at 1150°C for 12 hours. 131 for SiB3, CrB2, TaB2 and TiB2 compounds at 1423 K in equilibrium with the pure metal [150,151]. Figures 43 and 44 are the respective Si- and B-fluoride vapor pressures that were calculated for a pack comprised of Pure Si, MgF2 with either a CrB2 or TaB2 compound. The trend for the B- fluoride vapor pressures in Figs. 38, 40, 43 and 44 agrees with the boron activities in Table 25. A cross-section for the B-doped silicide coatings produced by packs composed of 7 wt.% Si, 2% MgF2 and A1203, with either 6% CrB2 or 6% TaB2 at 1150°C for 12 hours is shown in Figs. 45 and 46, respectively. An outer layer of TiB2 was detected by SEM/EDS and XRD (Appendix A, Figs. 147b and 147c), with a five-layer silicide coating revealed by etching. The thickest layer of TiB2 layer was observed for the B-doped silicide coating formed with a crB2 compound, in agreement with the higher B-fluoride vapor pressures. Chromium was not detected in the silicide or boride coating. However, the thick boride layer formed by the CrB2- containing pack was highly convoluted, not adherent to the silicide layer, and frequently peeled off the Ti-silicide during the cleaning procedure. The TiB2 layer formed on the silicide by a TaB2-containing pack was thinner, but more adherent to the silicide. A comparison of Figs. 41a, 45 and 46 shows that the formation of a thicker outer layer of TiB2 results in a thinner silicide coating. The TiB2 layers formed by packs 132 containing CrB2 or TaB2 were continuous, but quite porous. The resistance to the gas-phase diffusion of the Si- fluorides imposed by a thicker outer layer of TiB2 slows the growth kinetics for the silicide layers. The displacement reaction between TiSi2 and TiB2 (Eq. 4.18) probably occurs at a faster rate for a higher gas-phase flux of B-fluorides, which produces a thinner silicide coating. However, the morphology of the B-doped silicide coatings grown by packs containing CrB2 or TaB2 was an improvement over the undoped silicide coating formed by a MgF2 activator. Figure 47 is a plot of the vapor pressures for the Si- and B-fluorides calculated for a pack consisting of pure Si, a TiB2 compound and an A1F3 activator. A comparison of Figs. 40 and 47 reveals that the use of a less stable A1F3 activator (Table 11) increases the B-fluorides with respect to the Si-fluorides. A five-layered silicide coating with an outer layer of TiB2 is produced by a pack consisting of 7 wt.% Si, 6% TiB2, 2% A1F3, and A1203 at 1150°C for 12 hours, as shown in Fig. 48. The TiB2 layer is almost as thick as the layer produced by the pack comprised of 7 wt.% Si, 6% TaB2, 2% MgF2, and A1203 (Fig. 44), which has a higher boron activity. The TiB2 layer is continuous and porous. Therefore, the two ways of increasing the amount of boron in the B-doped silicide coating are: (1) to increase the chemical activity of the boron source, or (2) to use a less stable activator. 133 Figures 49, 50 and 51 are the calculated vapor pressures for the B- and Si-fluorides that result from packs composed of pure Si with either TaB2 and A1F3, TiB2 and CuF2, or TaB2 and CuF2. Higher vapor pressures for the B- fluorides with respect to the Si-fluorides are observed for a higher B activity, or for a less stable activator, in agreement with the previous results. The thickness of the TiB2 and TiSi2, TiSi, Ti5Si4, Ti5Si3 and Ti3Si layers that are produced by a pack comprised of pure Si, with either a TiB2 or TaB2 compound, and a MgF2, A1F3 or CuF2 activator and A1203 filler at 1150°C for 12 hours are summarized in Tables 26, 27, and 28, respectively. The thickness of the TiB2 layer increased by using a less stable activator or increasing the chemical activity for boron, in agreement with the trend predicted by the thermodynamic calculations. The most significant increase in the thickness of the outer TiB2 layer resulted from using TaB2 powder in the pack, which raised the activity of boron. The thickness of the silicide layers was smaller for a thicker outer layer of TiB2, which indicates that the thicker TiB2 layer slows the gas-phase diffusion of the Si-fluorides. Additionally, trace amounts of A1F3 or a-Al203 were detected at the surface of the coatings. Since CP-titanium undergoes an allotropic phase transformation from hexagonal a-phase to body-centered cubic 3-phase at 882°c (Fig. 2), the microstructure of the 134 E O £3 3V a. "5 ••e o a. Q t«00 Temperature (K) Figure 47. The equilibrium vapor pressures for the Si and B fluorides calculated for a pack comprised of pure Si, a TiB2 compound, and an A1F3 activator at 1000-1600 K. ^ ■ s o pm .1 ■ i ------il Figure 48. Optical micrograph of the polished and etched cross-section for a B-doped silicide coating formed by a pack comprised of 7 wt.% Si, 6% TiB2, 2% A1F3 and A1203 at 1150°C for 6 hours. 135 Table 26. The average thicknesses of the TiB2, TiSi2, TiSi, Ti5Si4, TisSi3 and Ti3Si layers that are produced by a pack comprised of either 7 wt.% Si, 6% TiB2, 2% MgF2 and A1203, or 7 wt.% Si, 6% TaB2, 2% MgF2 and A1203 at 1150°C for 12 hours. TiB2 [Mm] TaB2 [Mm] Containing Containing TiB2 1.21 6.20 TiSi2 46.4 35.6 TiSi 23.0 16.0 Ti5Si4 11.7 10.0 Ti5Si3 2.07 2.17 Ti3Si 3.11 3.17 s ir -3.0 £ 3 n § Q. "5 •■e o a. StF cn o. 100 1000 1100 1200 14001300 1000 Tem perature (K) Figure 49. The equilibrium partial pressures of the Si and B fluorides calculated for a pack comprised of pure Si, a TaB2 compound, and an A1F3 activator at 1000-1600 K. 136 Table 27. The average thicknesses ([Mm]) of the TiB2, TiSi2, TiSi, Ti5Si4, Ti5Si3 and Ti3Si layers grown by a pack comprised of either 7 wt.% Si, 6% TiB2, 2% A1F3 and A1203, or 7 wt.% Si, 6% TaB2, 2% A1F3 and A12o 3 at 1150°C for 12 hours. TiB2 [Mm] TaB2 [Mm] Containing Containing TiB2 2.39 7.76 TiSij 35.8 15.2 TiSi 21.4 24.1 Ti5Si4 10.7 12.6 TisSi3 1.83 2.96 Ti3Si 3.17 3.85 0J0 £o ¥ BF. 3 m w o> o *-e 0 Q. SIF 01 Q 1000 1100 1200 1300 1400 1900 1600 T em perature (K) Figure 50. The equilibrium partial pressures of the Si and B fluorides calculated for a pack comprised of pure Si, a TiB2 compound, and a CuF2 activator at 1000-1600 K. 137 Table 28. The average thicknesses {[fin]) of the TiB2, TiSi2, TiSi, Ti5Si4, Ti5Si3 and Ti3Si layers grown by a pack comprised of either 7 wt.% Si, 6% TiB2, 2% CuF2 and A1203, or 7 wt.% Si, 6% TaB2, 2% CuF2 and A12o3 at 1150"C for 12 hours. TiB2 [/im] TaB2 [jra] Containing Containing TiB, 2.78 7.89 TiSij 35.7 17.4 TiSi 21.1 24.1 Ti5Si4 11.6 11.8 Ti5Si3 2.61 2.44 Ti3Si 3.34 3.91 s StF. £ a n o *Eo 0. O 1000 1100 1400 1800 two Temperature (K) Figure 51. The equilibrium partial pressures of the Si and B fluorides calculated for a pack comprised of pure Si, a TaB2 compound, and a CuF2 activator at 1000-1600 K. CP-titanium workpiece is changed by coating at 1150°C for 12 hours. The microstructure changes that are produced by the diffusion coating anneal used for the aluminizing of Ni-base superalloys are reversed by heat treating the workpiece after the coating treatment, an approach that could be used for Ti-base alloys. To address this problem, a lower temperature and shorter time were used to determine whether thin coatings could provide the same protection as thick coatings. Based on the kinetic measurements, discussed in the proceeding section, a coating cycle of either 6 hr. or 28 hr. at 950®C was chosen to produce a coating that is 1/4 or 1/2 the thickness observed after 12 hours at 1150°C. The thicknesses of the coatings grown by a pack comprised of 7 wt.% Si, 6% TiB2, and A1203 with either a MgF2 or CuF2 activator are reported in Table 29. 4.4.2 Orowtb Kinetics of B-doped Silicide Coatings on CF-Titaniun The kinetics for the five-layer silicide coatings grown by packs composed of 7 wt.% Si, 6% TiB2 compound, and A1203 activator, with either 2% MgF2 or A1F3 activator were measured at 1150°C, 1050°C and 950°C. Since the TiB2 layer on a coating grown by an AlF3-activated pack was thick enough to be resolved, the growth kinetics for the TiB2 layer were measured. Figures 52 and 53 are the respective linear regression fits to a plot of thickness versus the square root of time for B-doped silicide coatings formed by 139 MgF2- and AlF3-activated packs at 1150°C. A good fit was observed for each layer, which indicates the kinetics are parabolic. The apparent rate constants were determined by accounting for the incubation time using Eq (4.14). The intrinsic rate constants for each five-layer silicide coating were determined from the apparent rate constants using Eq. (4.4), as reported in Tables 30 and 31. Figures 54 and 55 are the Arrhenius plots for the B-doped silicide coatings grown by a MgF2- and A1F3-activated pack, respectively. The activation energy for each layer of the B-doped silicide coatings are reported in Table 32. Figure 56a shows a parabolic fit for the growth of the TiB2 layer on the silicide coating formed by an A1F3- activated pack. To compare the measured growth kinetics for the two-layer boride coating on Ti with the single TiB2 layer on the silicide coating, a correction is made for the difference in the B activity gradient using Eq. (4.6), kP " kp[ra]{(lnaB[TiB2/TiB))/(lnaB[TiB2/Ti])) (4.19) where kpjmj is the parabolic rate constant measured from the single layer of TiB2, aB[TiB2/TiB] is the boron activity for equilibrium between TiB and TiB2, aB[TiB2/Ti) is the boron activity for equilibrium between TiB2 and Ti, and kp is the intrinsic rate constant used for comparison, which is reported in Table 31. The arrhenius plot for the intrinsic rate constant is shown in Fig. 56b, and the activation energy is reported in Table 32. 140 Table 29. The average thicknesses ([mw]) of the TiB2, Tisi2, TiSi, Ti5Si4# Ti5Si, and Ti3Sl layers grown by a pack composed of either 7 wt.% Si, 6 % TiB2, 2% MgF2 and A1203, or 7 wt.% Si, 6 % TiB2, 2% CuF, and AI 9O 3 at either 950°C for 6 hours, or 950°C l&r 28 hours. MgF2 [Mra] CuF2 [M®] Activator Activator 950°C, 6 hours TlB2 0.35 1.05 TiSl2 15.8 3.38 TlSl 4.98 2.15 Ti 5Si4 1.63 1.54 Tl 5Sl3 0.46 0.25 Ti3Si 0.79 0.49 950°C, 28 hours T1B2 0.52 1.22 TlSl2 27.4 24.2 TlSl 8.73 12.2 Ti 5Si4 2.38 2.89 Ti 5Si3 0.84 0.85 Ti3Si 1.44 1.14 141 Table 30. Intrinsic parabolic rate constants (kp[cm2/s]) measured for TiSi2, TiSi, Ti=Si4, Ti5si3, and Ti3Si layers grown on a B-doped silicide coating by a pack comprised of 7 wt.% Si, 6% TiB2, 2% MgF2 and A1203 at 950°C, 1050°C, and 1150°C. 1150°C 1050°C 950°C TiSi, 1.08X10-10 7.62X10"11 4.21X10-11 TiSi 1.05X10-10 2.80X10"11 9.21X10"12 Ti5Si4 4.81X10-11 7.50X10“12 1.84X10”12 T i5Si3 7.34X10-12 4.29X10**12 9.03X10“13 Ti3Si 6.49X10"12 2.84X10-12 2.44X10”12 £ £i 0 (/) W 0) £o / -- - - Jr.-/5 1- r tw T T I1)1 | I i i i p i i i^j i I i I | I 'l'l i | I i i i | ) t i I 50 100 150 200 250 300 350 Square Root Time [sec1/2] Figure 52. Linear regression fit to a plot of thickness versus time for TiSi2, TiSi, Ti5Si4, Ti5Si3, and Ti3Si layers formed on a B-doped silicide coating by a pack comprised of 7 wt.% Si, 6% TiB,, 2% MgF2 and A1203 at 1150#C. The error bars are one standard deviation. 142 Table 31. Intrinsic parabolic rate constants (kp(cm2/s]) determined for the TiB2 layer and the Tisi2, TiSi, Ti5Si4, Ti5Si3, and Ti3Si layers grown on a B- doped silicide coating by a pack comprised of 7 Wt.% Si, 6% TiB2, 2% AlFj and Al203 at 950°C, 1050°C, and 1150°C. 1150°C 1050°C 950°C TiSi2 8.62X10"11 5.49X10"11 5.40X10"12 TiSi 5.41X10"11 4.12X10-11 2.40X10"12 Ti5Si4 2.91X10"11 9. 52X10”12 1.15X10"12 Ti5Si3 2.78X10-12 2.66X10"12 9.52X10"13 Ti3Si 3.38X10-12 1.88X10"12 2.64X10"13 TiB2 5.31X10-14 3.48X10-14 1.01X10-14 (0 (0 a) 5 TiSi A O J. i 250 Square Root Time [se e n Figure 53. Linear regression fit to a plot of thickness versus time for Tisi2, Tisi, TisSi4, Ti5si3, and Ti3si layers formed on a B-doped silicide coating by a pack comprised of 7 wt.% Si, 6% TiB2, 2% A1F3 and A1203 at 1150°C. The error bars are one standard deviation. 143 Table 32. Activation energies (Q[J/mole]) determined from the intrinsic parabolic rate constants for the growth of B-doped silicide layers by packs composed of 7 wt%. Si, 6% TiB2 and A1203 filler with either MgF2 or A1F3 activator. MgF2-activated AlF3-activated TiSi2 68,500 203,000 TiSi 175,000 230,000 Ti5Si4 235,000 235,000 Ti5Si3 153,000 77,300 Ti3Si 70,200 162,000 TiB2 -- 121,000 1150C 1050C 950C 5 10‘i— ♦— ^—t ♦------r nb■ ■ r r ra «TiSi TiSi 8 lO-’S r d) ■ • TO 0 i0',2i r : ! 1*C C T j r i— I I I I 1 I | I I I 0.00065 0.0007 0.00075 0.0008 0.00085 0.0009 1/T TK'1! Figure 54. Linear regression fit to an Arrhenius plot of the intrinsic parabolic rates for the Tisi2, Tisi, Ti5si4, Ti5Si3, and Ti3Si layers formed by a MgF2-activated pack. Ti the intrinsic parabolic rates for the the for rates parabolic intrinsic the Figure 55. Linear regression fit to an Arrhenius plot of of plot Arrhenius an to fit regression Linear 55. Figure 5 Intrinsic Rate Constant [cm /s] Si3, 00065 0. 00075 0. 00085 0. 9 0 0 .0 0 5 8 0 0 .0 0 8 0 0 .0 0 5 7 0 0 .0 0 7 0 0 .0 0 5 6 0 0 .0 0 ii and and Ti Si Ti3Si 0 1050C 50C 1 1 layers formed by an AlF3-activated pack. AlF3-activated an by formed layers Ti Si / fK*1!1/T TiSi Ti.Si 950C TiSi2r Ti TiSi, TiSi 5 Si4, 144 145 1150C s' il § 0.8 - 1050C -i— i— i— ]— i— i i i | i i— i— i— i— i— i i 1 | ' 50 100 150 200 250 ,1/2i (a) Square Root Time [seclfii] 1150C 1050C 950C jn (M E u c *->ra t/y c o o 0) cc« 1 0 ' N o ’« c •c £ 10 (b) 0.00065 0.0007 0.00075 0.0008 0.00085 0.0009 1/T fK''l Figure 56. Linear regression plots for the TiB2 layer on the B-doped silicide coating grown by a pack comprised of 7 wt,% Si, 6% TiB2, 2% A1F3 and A1203; (a) thickness versus the square root of time, and (b) Arrhenius plot for the intrinsic rate constants. 146 The difference between the intrinsic rate constants for the undoped silicide (Table 14) and the B-doped silicide (Table 30) coatings grown by a MgF2 activator is less than a factor of 1.5, which indicates little difference in the magnitude of the growth kinetics. The activation energies for the growth of the B-doped silicide by a MgF2 activator are approximately the same as for the undoped silicide coatings. Therefore, the growth mechanism for the B-doped silicide is the same as the undoped silicide for a MgF2- activated pack, which is gas-phase diffusion controlled. The formation of the discontinuous TiB2 islands does not change the growth kinetics of the B-doped silicide coating. The difference between the intrinsic rate constants for the undoped (Table 15) and B-doped (Table 31) silicide coatings formed by an A1F3 activator is generally less than a factor of 2 to 4. The activation energies for the growth of each layer for the undoped (Table 17) and B-doped (Table 32) silicide coatings grown by an AlF3-activated pack are similar, except for the differences observed for the Ti5Si4 and TisSi3 layers. This result indicates that the growth mechanism for the B-doped silicide coating by an A1F3 activator is the same as the undoped silicide coating, which was solid-Btate diffusion. However, the inconsistencies in the kinetic data demonstrate that the formation of a nearly continuous TiB2 layer affected the growth of the silicide coating layers. The intrinsic rate constants for the TiB2 layer on the B-doped silicide coating (Table 31) are 1 to 2 orders in magnitude lower than the intrinsic rate constants measured for the growth of TiB2 from a pure boriding pack (Table 22). The activation energy for the growth of TiB2 on the silicide coating (Table 32) is similar to the activation energies determined from a boriding pack (Table 24). These results indicate that the growth mechanism for the TiB2 layer on the silicide coating is approximately the same as for the formation of a TiB2 diffusion coating, which was assumed to be solid-state diffusion in the bulk. The porosity and slower kinetics for the growth of the TiB2 layer indicate that a surface step may affect the kinetics. 4.4.3 Boron-Doped Silicide Coatings on Ti-Al-Nb Alloys The successful pack processes that were developed for CP-titanium were also used to coat Ti-aluminide alloys. Figure 57a is the polished and etched cross-section of a B- doped silicide coating grown on a Ti-22Al-27Nb alloy by a pack composed of 7 wt.% Si, 6% TiB2, 2% MgF2 and A1203 after 12 hours at 1150°C. The dual-layer coating consists of a thick outer layer of TiSi2 and a thin TiSi layer adjacent to the substrate. The EDS profile in Fig.' 67b shows that aluminum is concentrated at the surface, but minor amounts (< 0.5 at.%) of Al are present within the bulk of the TiSi2 and TiSi layers, with an Al enrichment observed at the coating/alloy interface. Niobium is present in the silicide 148 coating, and is depleted in the interdiffusion zone adjacent to the coating/alloy interface. The interdiffusion zone is highlighted upon etching, and is the same width as the Al enrichment zone in Fig. 57b. Ti-Al-Nb alloys with more than 30 at% Al possess better high-temperature oxidation resistance than orthorhombic-base alloys [22]. The Al- enriched interdiffusion zone may serve as a second line of defense in the event of failure of the silicide coating. X-ray diffraction of the coating surface detected TiSi2 (JCPDS# 10-225, S.G. #63) with NbSl2 (JCPDS# 8-450) and minor peaks of TiB2 and TiAl3 (Appendix A, Fig. 148a). TiSi2 is known to undergo a polymorphic phase transformation from a metastable base-centered orthorhombic phase (C49) to the stable face-centered orthorhombic phase (C54) at high temperature [153], The polymorph for the TiSi2 coating on the Ti-22Al-27Nb alloy is the metastable C49 phase, but the TiSi2 layer grown on CP-titanium was the stable C54 phase. Previous work has shown that Al stabilizes the metastable C49 TiSi2 phase, which indicates that small amounts of Al are included in the coating [154]. However, the solubility for Al in TiSi2 is very low [ 154 ], so that minor amounts of Al were incorporated into the Ti-silicide layers, and the diffusive conversion of the alloy into the Ti-silicide coating results in an enrichment of Al at the coating/alloy interface. 149 The backscattered electron (BSE) SEM image of the coating in Fig. 58a reveals bright regions that are rich in Nb. Since NbSi2 was detected by XRD, these results demonstrate that NbSi2 precipitates formed in the TiSi2 coating, and that TiSi2 has some solubility for Nb at high temperature. Thus, Nb from the alloy was dissolved into the growing Ti-silicide layers at high temperature, and formed Nb-silicide precipitates, which explains the Nb depletion in the interdiffusion zone. The microprobe trace in Fig. 58b shows that all the boron is concentrated in a 1 p layer at the coating surface, in agreement with the results observed for the B- doped silicide coating on CP-titanium. Only trace amounts of condensed activator salt were detected at the surface of these coatings, similar results were obtained for the growth of B-doped silicide coatings on the Ti-20Al-22Nb, Ti- 22A1—23Nb, and Ti-24Al-llNb alloys. The thickness of the B-doped silicide coating grown on the Ti-22Al-27Nb alloy (Fig. 57a) is much greater than the same coating formed on CP-titanium (Fig. 41a) by the same HAPC method. The dissolution of Nb and Al during the formation of the coating may result in faster growth kinetics. One problem associated with the faster growth kinetics for the alloy substrates is convolution formation at the corners, as shown in Fig. 59 for the B-doped silicide coating on a Ti-22Al-27Nb alloy. The volume change v v. • / >' Substrate fv I 100pm ^ 100 90- • TiSi2 & -TiSi Ti—22AI—27Nb 80- - NbSi2 Substrate E 70- ■Hd—cl 1 60- Silicon Profile c o 50- j3 *5i o 40- Niobium Profile CL . A * -o- E 30- ^A<0-0 o — ,Q~ -O- - 0- o '"O * A- -A- -A- o 20- -A- Alumtnum Profile 10- 0- 100 200 300 400 500 600 Distance From Outer Edge of Coating (jim) (b) Figure 57. B-doped silicide coating grown on Ti-22Al-27Nb by a pack comprised of 7 wt.t si, 6% TiB2, 2% MgF2 and A1203 at 1150°C for 12 hours, (a) Optical micrograph of the polished and etched cross-section and (b) EDS profile of si, Nb and Al in the coating. 151 " n m n s ^ - NbSIj ppt. ESI Baat-Alloy *nSi2+NbSi2 TiSi - Substrote K E Silicon Profile o ced - □—□— D- < c Aluminum o JO Profile 'w o a <- Boron Profile E o o Niobium Profile O 0 Figure 58, B-doped silicide coating grown on Ti-22Al-27Nb by a pack composed of 7 wt.% si, 6% TiB2, 2% MgF2 and Al2o3 at 1150• C for 12 hours, (a) SEM/BSE micrograph of a polished cross-section and (b) microprobe profile of B, Si, Nb and Al in the coating. 152 Figure 59. Optical micrograph of the corner of a B-doped silicide coating grown on a Ti-22Al-27Nb alloy by a pack comprised of 7 wt.% si, 6% TiB2, 2% MgF2 and A1203 at 1150°c for 12 hours, which shows the convolution. associated with the diffusive conversion of the substrate into the TiSi2 coating probably produced this effect. Additionally, the Ti-22Al-27Nb alloy is cycled through the beta transus temperature by coating at 1150°C, which may have produced the distortions at the corners. One solution for this problem is to use a lower temperature and shorter time to produce a thinner coating. Figure 60a is a SEM micrograph of a B-doped silicide coating grown on a Ti-20Al-22Nb alloy by a pack comprised of 7 wt.% Si, 6% TiB2, 2% MgF2 and A1203 at 950°c for 6 hours. The coating consists of a thick outer layer of TiSi2 with a thin layer of TiSi adjacent to the substrate. The EDS profile in Fig. 60b shows that the concentration of Nb in the coating is similar to that observed for the thicker 153 coating, and B is concentrated at the coating surface. The phases detected by X-ray diffraction are the same as those observed on the thicker coating, and the TiSi2 is the metastable C49 phase (Appendix A, Fig. 148b) One big difference between the thin and thick coatings is the narrow interdiffusion zone at the coating/alloy interface in Fig. 60a. The Al concentration in the interdiffusion zone of the thin coating is 50 at.%, which indicates this was a 2 to 4 (in layer of y-TiAl. The y-TiAl layer possess good oxidation resistance and may serve as a second barrier against the penetration of oxygen following any failure of the silicide coating. Large changes were observed in the microstructure of the base alloy coated at 1150°C for 12 hours in Fig. 57a, but little change was observed for the alloy coated at 950°C for 6 hours in Fig. 60a. Thus, one advantage of coating the Ti-20Al-22Nb and Ti-22Al-27Nb alloys at 950°C is that the beta transus temperature was not exceeded. The interdiffusion coefficients for the beta phase (bcc) are much higher, which results in a wide interdiffusion zone and less Al enrichment (Fig. 57b). Coating the alloy below the transformation temperature results in a narrow interdiffusion zone, and greater Al enrichment (Fig. 60b). The successful endeavors for increasing the B concentration of the B-doped silicide coatings on CP- titanium directed new pack coating treatments for the 154 — * -si 100 TiSij Coating TiSi Substrate 80- o Boron Si Profile •g 2 60- A1 Profile •oo 0e •a VI 40- 1 3 2 0 - Nb Profile (b) d(um) Figure 60. B-doped silicide coating grown on Ti-20Al-22Nb by a pack composed of 7 wt.% si, 6% TiB-j, 2% MgF, and Al2o3 at 950°C for 6 hours, (a) SEM micrograph of a polished and etched cross-section and (b) EDS profile of B, Si, Nb and Al in the coating. 155 Ti-Al-Nb alloys. For the Ti-20Al-22Nb alloy these were composed of 7 wt.% Si, 6% TaB2, 2% MgF2, and A1203 at 1150°C for 12 hrs., 7 % Si, 6% TiB2, 2% CuF2, and A1203 at 950°C for 6 and 28 hrs., and 7% Si, 6% TaB2, 2% CuF2, and A1203 at 950°c for 6 and 28 hours. The average thicknesses of the TiB2, TiSi2 and TiSi layers measured for all the B-doped silicide coatings on the Ti-Al-Nb alloys used in this study are summarized in Table 33. Increasing the boron activity in the pack or using a less stable activator resulted in thicker layers of TiB2 at the coating surface and thinner silicide layers, in agreement with the results for CP- titanium. Additionally, the thickness of the coating layers grown by a pack treatment composed of 7 wt.% Si, 6% TiB2, 2% MgF2, and A1203 at 1150°C for 12 hours on the Ti-22Al-27Nb and Ti-20Al-22Nb alloys were similar. The maximum Al enrichment in the interdiffusion zone at the coating/alloy interface was 30-35 at% for coatings grown at 1150°C for 12 hours, and 49-54 at% for a coating cycle at 950°C for 6 or 28 hours, with a thicker zone observed for the coating grown at 950°c for 28 hours. The same amount of Nb was always observed in the thick and thin coatings, and the same phases were detected by XRD, with larger amounts or A1F3 and a-Al203 detected at the surface of coatings grown using the less stable CuF2 activator. Figure 61 shows the thick, porous coating that was formed on a Y"TiAl Ti-48Al-2Mn-2Nb alloy by a pack composed 156 of 7 wt.% Si, 6% TiB2, 2% MgF2, and Al203 filler at 1150°C for 12 hours. This coating was not protective, and changing the activator, boride compound, filler, coating time and temperature did not correct this problem. Silicon is probably deposited from the fluoride vapor by a displacement reaction (Eg. 1.3) with the Al contained in TiAl to produce a more stable A1F3, (3/x) SiFx (v) + Al(in TiAl) = Si(in TiSi2) + AlF3(g) (4.20) The high activity for Al in TiAl results in this displacement mechanism that deposits Si by extracting Al from the workpiece to produce a porous silicide coating. The activity for Ti in the TiAl compound is low, which does not allow the exclusive growth of a TiSi2 coating by the HAPC method. 4.5 Germanium-Doped silicide Coatings Only fluoride-base activators were used to produce the Ge-doped silicide coatings in the study. Successful results were never obtained for a chloride-base activator, as with the case of the B-doped and undoped silicide coatings. The approach for this study was adopted from the work of Mueller et al. [25-27] for producing Ge-doped MoSi2 coatings. However, a good coating was not achieved using the same process developed by Mueller at al. [25-27], and further work was required to optimize the coating process. 157 Table 33. The average thickness < Cixm]) for the TiB2, Tisi2 and TiSi layers for B-doped silicide coatings grown on Ti-22Al-27Nb (22-27) and Ti-20Al-22Nb (20-22) by 8 different pack treatments; (1) Si-TiB2/MgF2/Al203 at 1150°C for 12 hours, (2) Si-TiB2/MgF2/Al203 at 950#C for 28 hours, (3) Si-TiB2/MgF2/Al203 at 950°C for 6 hours, (4) Si-TaB2/MgF2/Al203 at 1150°C for 12 hours, (5) Si-TiB2/CuF2/Al203 at 950°C for 28 hours, (6) Si-TiB2/CuF2/Al203 at 950#C for 6 hours, (7) Si-TaB2/CuF2/Al203 at 950°C for 28 hours, (8) Si-TaB2/CuF2/Al203 at 950°C for 6 hours,. TiSij TiSi TiB2 (1) [22-27] 102.6 10.44 0.95 (1) [20-22] 113.3 12.85 0.92 (2) [20-22] 41.20 5.76 0.82 (3) [20-22] 17.74 3.78 0.53 (4) [20-22] 76.74 16.18 2.08 (5) [20-22] 37.08 4.60 1.29 (6) [20-22] 9.07 2.30 0.87 (7) [20-22] 17.18 4.11 2.02 [20-22] 7.24 2.15 1.48 4.5.1 Ge-Doped Silicide coatings on CP-Titanium The calculated vapor pressures for the Si- and Ge- fluorides that result from a pack comprised of pure Si and Ge, a MgF2 activator and Sic filler are shown in Fig. 62. The partial pressures for the Ge-fluorides are only slightly lower than the Si-fluorides, which indicates that a silicide coating with Ge additions should be formed by this pack. Figure 63 shows that a five-layered Ge-doped silicide coating was produced by a pack composed of 12 wt.% Si, 6% Ge, 2% MgF2 and Sic. No second phases are observed in any of the layers, and the Ge solute was not uniformly distributed in the five- layered coating. The only phase at the surface of the coating 158 Figure 61. Optical micrograph of the polished cross-section of a B-doped silicide coating grown on a Ti-48Al-2Mn-2Nb by a pack comprised of 7 wt%. si, 6% TiB2, 2% MgF2, and A1203 at 1150<*C for 12 hours. ■ S .0 a n *)m SJF a '■£o CL SF cno •A 1000 1*00 1900 1400 Tem perature (K) Figure 62. Thermodynamic calculation of the Si and Ge fluorides for a pack composed of si, Ge, MgF2 and Sic at 1000- 1600 K. 159 Eanm t . m |T I(8l,Q i) G r u m p y 100 90 TiSIf TiSi TleS] -*n3si 80 E 70 T'ss h I 60 O 00 OCKB- c o 50 Silicon Profile \ |S *5» Ti o 40 Q. ■ x Base E 30 o a 20 10.. Germanium Profile to DO 0 - — r T — I l---1---h- y g i ¥ H h ft + 10 20 30 40 50 60 70 80 90 100 110 120 130 140 (b) Distance From Outer Edge of Coating (/urn) Figure 63. Ge-doped silicide coating formed by a pack comprised of 12 wt.% Si, 6% Ge, 2% HgF2 and Sic at 1150°C for 12 hours, (a) Optical micrograph and (b) EDS profile for Si and Ge. 160 detected by XRD was Tisi2 (JCPDS# 35-785, S.G. #70), with no condensed activator detected (Appendix A, Fig. 149a). The maximum Ge concentration in the Ti(Si,Ge)2 layer was observed at the Ti(Si,Ge)2/Ti(Si,Ge) interphase boundary with the minimum at the Ti(Si,Ge)2/gas interface. TiSi2 and TiGe2 have the same crystal structure, and form an isomorphous Ti(Si,Ge)2 solid solution over the complete range of composition [57]. The Ge concentration gradient in the Ti(Si,Ge)2 layer is not large, and can be explained by chemical demixing in the chemical potential gradient for si [103]. If the diffusivity for Ge is faster than for Si in the Ti(Si,Ge)2 solid solution, then the Ge would be expected to segregate to the Ti(Si,Ge)2/Ti(Si,Ge) interface for uncorrelated, random diffusion. Gas et al. [151] measured the Ge diffusion coefficient in TiSi2 for ion-implanted Ge in a TiSi2 thin film on a Si wafer, and determined that the diffusion coefficient for Ge in TiSi2 is 2 to 3 orders of magnitude slower than for Si in TiSi2. If diffusion of Ge in Ti(Si,Ge)2 is slower than Si, then the experimentally observed Ge concentration profile should be reversed. The Ge diffusion rates determined by Gas et al. [151] are tainted by the effects of ion implantation (increased vacancy concentration), and the Ge concentrations are much smaller than those used in this study. Since the TiSi2 layer used by Gas et al. [151] was on a pure Si wafer and no higher silicide phase can form, the chemical potential gradient that drives the interdiffusion 161 of Ge in TiSi2 is different than the driving force provided by a Ti substrate. Therefore, the results of Gas et al. [151] are not applicable to the growth of Ge-doped Ti-silicide diffusion coatings. The solid-state diffusion rate for the Ge solute probably exceeds Si for the Ti(Si,Ge)2 solution, and the Ge concentration profile is explained by simple demixing in a chemical potential gradient. However, the interaction coefficients (Lsi#Ce and boe,si) * which describe the influence of the Si concentration gradient on the solid-state diffusion of Ge and vise-versa, are required for a quantitative description of chemical demixing in Ti(Si,Ge)2. The Ge concentration in the Ti(Si,Ge) phase is uniform, but lower than Ti(Si,Ge)2. The Ge composition in each layer is the same when expressed in terms of the mole fraction of TiGe2 in Ti(Si,Ge)2, and TiGe in Ti(Si,Ge). The Ti(Si,Ge) layer in Fig. 63a is significantly thinner than the undoped TiSi layer grown by a MgF2-activated pack in Fig. 22a. A TiGe phase does not exist in the Ti-Ge system [57], and the Ge additions may serve slow the growth rate of the Ti(Si,Ge) layer. The growth kinetics for the Ge-doped Ti-silicides are discussed in the proceeding section. The highest concentration of Ge in the five-layered silicide coating was observed in the Ti5(Si,Ge)4 layer, and the maximum Ge concentration within the Ti5(Si,Ge)4 layer was observed at the Ti5(Si,Ge)4/Ti5(Si,Ge)3 interface. The diffusion of Ge in the Ti5(Si,Ge)4 solid solution is probably 162 faster than Si, which results in demixing in the chemical potential gradient and the enrichment of Ge at the Ti5(Si,Ge)4/Ti5(Si,Ge)3 interphase boundary [103]. Additionally, the diffusion of Ge may be faster than Si in the Ti(Si,Ge)2, Ti(Si,Ge) and Ti5(Si,Ge)4 layers, which results in global chemical demixing in the chemical potential gradient through the three layers and explains the enrichment of Ge in the Ti5(Si,Ge)4 phase. The Ti-germanide phase that is closest to the composition of Ti5Si4 is the Ti5Ge6 phase, which may dissolve into Ti5Si4. However, multi-layered growth kinetics complicate the expectations for chemical demixing, which are written in terms of diffusion in a single layer [103]. The Ge-concentration is lower in the thin Ti5(Si,Ge)3 and Ti3(Si,Ge) layers adjacent to the coating/metal interface. The Ti5Ge3 and Ti3Ge phases form isomorphous solid solutions with Ti5Si3 and Ti3Si. The lower Ge concentration in the Ti- rich Ti5(Si,Ge)3 and Ti3(Si,Ge) layers is not a true decrease in Ge, because the mole fractions of Ti5Ge3 and Ti3Ge in the Tis(Si,Ge)3 and Ti3(Si,Ge) solid solutions is similar to those for Ti(Si,Ge)2, Ti(Si,Ge) and Ti5(Si,Ge)4. Chemical demixing may occur in Ti5(Si,Ge)3 and Ti3(Si,Ge), but these inner layers were too thin to resolve the Ge profile within each layer. The vapor pressures for the Si- and Ge-fluorides that were calculated for a pack composed of pure Si, pure Ge and either an A1F3 or CuF2 activator are shown in Fig. 64. The 163 magnitudes of the Si and Ge fluorides are higher, and the Ge fluorides increase relative to the Si fluorides for a less stable activator. Therefore, more Ge-doping of the Ti- silicides should result for a less-stable activator. However, the coatings produced by packs comprised of pure Si, pure Ge, and a Sic filler with either an A1F3 or CuF2 activator were non-adherent and highly convoluted. The reason for these failures was not apparent. Figures 65a and 65b show that Ge-doped silicide coatings were achieved for packs composed of either 16 wt.% Si, 8% Ge, 2% A1F3 and A1203, or 16% Si, 8% Ge, 2% CuF2 and A1203, respectively. The morphology for the five-layered silicide coatings is generally the same as the coating grown by a MgF2- activated pack, i.e. the Ti(Si,Ge)2 and Ti5(Si,Ge)4 layers are thick, with thin Ti(Si,Ge), Ti5(Si,Ge)3, and Ti3(Si,Ge) layers. However, the thicknesses of the silicide layers are smaller, and the surface of the outer Ti(Si,Ge)2 layer is irregular with inclusions and a deposit at the surface. The inclusions were determined by EDS and XRD to be a salt-oxide deposit composed of A1F3 and cr-Al203, and the outer TiSi2 is the stable C54 phase (JCPDS# 35-785, S.G. #70), (Appendix A, Fig. 149b). Figure 66 shows that these coatings were convoluted at sharp corners. Table 2 shows that many variations in pack chemistry were attempted to eliminate the convolutions, but none were successful. The pack compositions that are listed in Table 4 minimized the convolutions and salt 164 E O 2 3 nV> o> o r o Cl. cn o —10JJ 1000 1100 1200 1300 1400 1500 1600 (a) Temperature (K) 0.0 E o 3.0 S1F. 3 n n SJF. P a ■•e o C«F a. SIF ot C4 O 9J0 1000 1100 1200 1300 1400 1500 1600 (b) Temperature (K) Figure 64. Thermodynamic calculation of the partial pressures of the Si and Ge fluorides for a pack comprised of (a) Si. Ge# and A1F3 and (b) si, Ge, and CuF2 at 1000-1600 K. 165 Figure 65. Ge-doped silicide coating (a) Optical micrograph for the polished cross-section of coating formed by a pack comprised of 16 wt.% Si, 8 % Ge, 2% A1F3 and AI 2O3 at 1150°C for 12 hours, and (b) SEM micrograph of a coating grown by a pack composed of 16% Si, 8 % Ge, 2% CuF2 and A120 3 at 1150°C for 12 hours. 166 looum Figure 66. A convolution in a Ge-doped silicide coating grown at the corner of a CP-titanium workpiece by a pack composed of 16 wt.% Si, 8% Ge, 2% CuF2 and A1203 at 1150°C for 12 hours. inclusions. Additionally, the least convoluting was observed for a CuF2-activated pack. Coatings that were produced by a pack consisting of pure Si and Ge, a MgF2 activator and A1203 filler were more convoluted than the A1F3- and cuF-activated packs, and were not studied in any detail. The composition profiles for the Ge-doped Ti-silicide coatings grown by the A1F3- and CuF2-activated packs are shown in Fig. 67. The amount of Ge observed in each of these coatings is less than that observed for the MgF2-activated pack in Fig. 63b; Nevertheless, the trends observed for the Ge concentration profiles in Fig. 67 are similar to those observed in the coating grown by a MgF2-activated pack. The Ge concentration in the coating grown by a CuF-activated pack (Fig. 67b) is larger than the AlF3-activated pack (Fig. 67a), 167 in agreement with the thermodynamic calculations of Fig. 64. Packs composed of 16 wt.% Si, 8% Ge, and A1203 filler with either 2% A1F3 or 2% CuF2 activator, and 12% Si, 6% Ge, 2% MgF2 and Sic were treated at 950°C for either 6 or 28 hours to produce a coating 1/4 or 1/2 the thickness of those grown at 1150°C for 12 hours. The coatings produced by packs comprised of Si-Ge/A1F3/Al203 and Si-Ge/CuF2/Al203 were highly convoluted. A coating cycle of 950°C for 12 hours was required to produce coatings that were adherent enough to survive a mild cleaning treatment. The thicknesses of the coatings resulting from packs comprised of Si-Ge/MgF2/SiC, Si- Ge/A1F3/Al203 and Si-Ge/cuF2/Al203 for the three coating cycles are summarized in Tables 34, 35, and 36, respectively. The undoped silicide coatings grown from packs composed of either Si/A1F3/Al203 or Si/CuF2/Al203 were adherent and not convoluted (Figs. 22b and 22c). However, convolutions were observed for Ge-doped silicide coatings grown by packs comprised of either Si-Ge/A1F3/Al203 or Si-Ge/CuF2/Al203. The molar volumes for the Ti-germanides are much larger than the Ti-silicides, and the volume change resulting from the diffusive conversion of the Ti workpiece into the Ge-doped silicide coating may produce the convolutions. The addition of Ge to the Ti-silicide may change the diffusion mechanism for the growth of the coating and produce convolutions. Inclusions were observed at the surface of the outer Ti(Si,Ge)2 layer for coatings grown by A1F3- and 168 100 90- TiSir TiSi •n5si4 100 Tipi.Ge)* Tis(Si,Ge)4 Ti Base 80-: Ti(SI,Ge)- Metal Ti}(SI,Gc)- Si-Profile ,t a . Tij(Si,Ge)* a 60-i 0 S 40-j k s1 8 20 H Ge-Profile 0 Tmm niiiTnni|iiiinihnrnniii|iimmi|imm iT]iniiilinlniiiiiii|iniTTm 20 40 60 80 (b) Distance [|im] Figure 67. EDS profiles of Si, Ge and Al for theGe-doped silicide coating (a) formed by a pack composed of 16 wt.% Si, 8% Ge, 2% A1F3 and A1203 at 1150«C for 12 hours, and (b) formed by a pack composed of 16 wt.% si, 8% Ge, 2% CuF2 and A1203 at 1150®C for 12 hours. 169 Table 34. The average thicknesses {[/Jm]) for the Ti(Si,Ge)2, Ti(Si,Ge), Ti5(Si,Ge)4, Tis(Si,Ge)3, and Ti3(Si,Ge) layers grown on CP-titanium by a pack comprised of 12 wt.% Si, 6% Ge, 2% MgF2, and Sic for three coating treatments: (1) 1150°C for 12 hours (2) 950°C for 28 hours (3) 950°C for 6 hours 1150°C 950°C 950°C 12 hours 28 hours 6 hours Ti(Si,Ge )2 65.30 21.37 8.44 Ti(Si,Ge) 11.20 4.54 2.35 Ti5(Si,Ge)4 32.10 11.34 7.33 Ti5(Si,Ge)3 2.20 1.08 0.69 Ti3{Si,Ge) 3.35 1.71 1.09 Table 35. The average thicknesses ([pm]) for the Ti(Si,Ge)2, Ti(Si,Ge), Ti5(Si,Ge)4, Ti5(Si,Ge)3, and Ti3(Si,Ge) layers grown on CP-titanium by a pack comprised of 16 wt.% Si, 8% Ge, 2% A1F3, and A1203 for three coating treatments: (X) 1150°C for“ 12; hours (2) 950°c for 28 hours (3) 950°C for 12 hours 1150°C 950°C 950°C 12 hours 28 hours 12 hours Ti(Si,Ge)2 36.30 22.56 7.37 Ti(Si,Ge) 15.29 5.67 1.72 Ti5(Si,Ge)4 21.97 6.19 2.59 Ti5(Si,Ge)3 2.79 0.83 0.66 Ti3 (Si,Ge) 3.09 1.39 0.96 170 Table 36. The average thicknesses ([/xm]) for the Ti(Si,Ge)2, Ti(Si,Ge), Tis(Si,Ge)4 , Tis(Si,Ge)3, and Ti3 (Si,Ge) layers grown on CP-titanium by a pack comprised of 16 wt.% Sir 8 % Ge, 2 % CuF2, and Al20 3 for three coating treatments: {1) 1150°C for 12 hours (2) 950°C for 28 hours (3) 950°C for 12 hours 1150°C 950°C 950°C 12 hours 28 hours 12 hours Ti(Si,Ge)2 24.83 18.40 12.14 Ti(Si,Ge) 5.47 5.67 1.72 Tig(Si,Ge) 4 33.46 10.74 5.39 T i 5 (Si,Ge)3 3.59 1.03 0.60 T i 3 (Si,Ge) 4.46 1.51 1.14 CuF2-activated packs, which should not occur for an diffusion mechanism. However, the salt inclusions were not observed at the inner layers, which are expected for a true outward diffusion mechanism. Additionally, convolutions were not observed for the Ge-doped silicide coating grown by a MgF2-activated pack with a Sic filler. The condensation of a salt layer or a complex deposition step at the surface that results from using an A1203 filler might produce the convolutions. The thicknesses of Ge-doped silicide coatings grown by the A1F3- and CuF2-activated packs were non-uniform, but uniform thicknesses were observed for the growth of the undoped silicide coatings, which indicates that gas-phase diffusion or a surface reaction could be the rate-limiting step. 171 4.5.2 Variation of Ge-content Mueller et al. [25-27] showed that the amount of Ge in the Mo-siliclde coating can be modified by the ratio of Si and Ge in the powder pack. This approach was used to change the Ge content in the Ge-doped silicide coatings grown on CP- titanium (Table 5). The coating cycle of 12 hours at 1150°c was used for all the pack compositions in Table 5. Figures 68/ 69 and 70 show that the Ge content is the highest for a low ratio of Si to Ge in the pack. However, most convolutions were observed on the coatings grown by packs with low Si to Ge ratios. Convolutions were observed for the pack comprised of 1:2 Si-Ge/MgF2/SiC, while no convolutions resulted from a pack composed of 2:1 Si-Ge/MgF2/SiC. Since a large salt-oxide deposit was not observed at the surface of the coatings grown by a pack composed of Si-Ge/MgF2/SiC, this result demonstrates that the addition of excessive Ge produced the convolutions. The molar volumes for the Ti-germanides are larger than the Ti-silicides, and the volume change that results from the diffusive conversion of the Ti into the Ge-doped silicide coating may have caused the convolutions. In thiB case, larger amounts of Ge in the silicide coatings would produce greater volume changes, and increase the number of convolutions. The average thicknesses for the Ge-doped silicide coatings grown by the three packs with the three different ratios of Si to Ge are reported in Tables 37, 38, and 39. A 172 thinner Ge-doped silicide coating was produced for the sane coating cycle by a pack with a lower Si to Ge ratio, which contains higher anounts of Ge within the silicide coating. The larger convolutions that were produced by the addition of nore Ge also resulted in thinner coatings. 4.5.3 Growth Kinetics of Gs-Dopsd silicide Coatings The growth kinetics for the Ge-doped silicide coatings grown by packs conposed of: (1) 12 wt.% Si, 6% Ge, 2% MgF2., Sic, and (2) 16% Si, 8% Ge, 2% A1F3 and A1203 were studied. A linear regression fit for each of the Ge-doped silicide layers is shown in Fig. 71 for the coatings grown by a MgF2- and AlF3-activated pack. The apparent rate constants were determined by correcting for the incubation time using Eq. (4.14), and Eqs. (4.4a) to (4.4e) were used to determine the intrinsic rate constants from the apparent rate constants. The molar volume (VA) for the Ge-doped titanium-silicides, which is required to determine the intrinsic rate constants by Eqs. (4.4), was determined by a rule of mixtures relationship with the assumption that the Ge concentration in each layer is constant: V Tiv{Si,Ge) B {XrivGe) V[TiyGe] + U-XTivGemTivSi] * (4-21) where VTiv^si#Caj is the molar volume of the Ge-doped titanium- silicide, XTivGo is the mole fraction for the germanide phase calculated from the measured Ge concentration, V[TivSi] is the molar volume for the silicide phase and V[TivGe] is the molar 173 Table 37. The average thicknesses ([/m]) for the Ti(Si,Ge)2, Ti(Si,Ge), Tis(Si,Ge)4, Ti5(Si,Ge)3, and Ti3(Si,Ge) layers grown on CP-titanium by a pack comprised of Si, Ge, 2% MgF2, and Sic at 1150*C with three Si to (1) 12 wt!% Si and 6% Ge [2:1] (2) 12% Si and 12% Ge [1:1] (3) 6% Si and 12% Ge [1:2] [2:1] [1:11 [1:2] Ti(Si,Ge)2 65.30 52.56 45.86 Ti(Si,Ge) 11.20 7.53 6.59 Tis(Si,Ge)4 32.10 31.64 32.69 Tl5(Si,Ge)3 2.20 3.01 2.53 Ti3(Si,Ge) 3.35 3.90 4.06 1:2 Pack Base: 1:1 Pack Meta L u cq oC 3 8. 4-: e S 2:1 Pack |iiii|im|Mi!jmi|iriT|mi|mi|iiii|iiM|iin|ii!i|im|im|iin 0 20 40 60 100 120 140 Distance [pm] Figure 68. EDS profile of Ge for the Ge-doped silicide coating formed by a pack comprised of Si, Ge, 2 wt.% MgF2 and Sic at 1150°C for 12 hours with three ratios of Si to Ge: (1) 12% Si and 6% Ge [2:1], (2) 12% Si and 12% Ge [1:1], and (3) 6% Si and 12% Ge [1:2]. 174 Table 38. The average thicknesses ([/xn]) the Ti(Si,Ge)2, Ti(Si,Ge), Tis(Si,Ge)4, Ti5(Si,Ge)3, and Ti3(Si,Ge) layers grown on CP-titanium by a pack comprised of Si, Ge, 2 wt.% A1F3, and A1203 at 1150°C for 12 hours with three Si to Ge ratios: (1) 16 Wt.% Si and 8% Ge [2:1] (2) 16% Si and 16% Ge [1:1] (3) 8% Si and 16% Ge [1:2] [2:1] [1:1] [1:2] Ti(Si,Ge)2 36.30 23.95 22.06 Ti(Si,Ge) 15.29 8.21 8.42 Ti5(Si,Ge)4 21.97 23.63 22.62 Ti5(Si,Ge)3 2.79 2.51 2.43 Ti3(Si,Ge) 3.09 4.06 3.70 1:1 Pack 8-^ 1:2 Pack Base Metal . e g , o * a 6 i B* o ;a *55 a 4-; B 8 a o 2*; 2:1 Pack II|IIII1IIII|I1I1III II|IIIM1III|TI1IIMII|IT1IIIIII|IIIIIIIII|IIIIIIIII 0 20 40 60 80 Distance Qim] Figure 69. EDS profile of Ge for the Ge-doped silicide coating formed by a pack comprised of Si, Ge, 2 wt.% A1F3 and A1,03 at 1150°C for 12 hours with three ratios of Si to Ge: (1} 16% Si and 8% Ge [2:1], (2) 16% Si and 16% Ge [1:1], and (3} 8% Si and 16% Ge [1:2]. 175 Table 39. The average thicknesses {[/xm]) for the Ti(Si,Ge)2, Ti(Si,Ge), Tis(Si,Ge)4, Tis(Si,Ge)3, and Ti3(Si,Ge) layers grown on CP-titanium by a pack comprised of Si, Ge, 2 wt.% Cu F2, and A1203 at 1150°C for 12 hours with three Si to Ge ratios: (1) 16 wt.% Si and 8% Ge [2:1] (2) 16% Si and 16% Ge [1:1] (3) 8% Si and 16% Ge [1:2] [2:1] [1:1] [1:2] Ti(Si,Ge )2 24.83 24.44 22.13 Ti(Si,Ge) 5.47 8.90 7.87 Tis(Si,Ge)4 33.46 30.04 34.35 Ti5(Si,Ge)3 3.59 2.62 2.33 Ti3 (Si,Ge) 4.46 4.24 4.43 12 1:2 Pack .rt. c 9-E a 1:1 Pack ;a *55 o a. 6- Base o6 Metal U 2:1 Pack 'niniiiniiinmniiin)iin)iiiiiiinniiiiiin)iiiiinmiimHmiiiiiiu|imiiiff ) 20 40 60 80 Distance Qun] Figure 70. EDS profile of Ge for the Ge-doped silicide coating formed by a pack comprised of Si, Ge, 2 wt.% CuF2 and A1203 at 1150°C for 12 hours with three ratios ofSi toGe: (1) 16% Si and 8% Ge [2:1], (2) 16% Si and 16% Ge [1:1], and (3) 8% Si and 16% Ge [1:2]. 176 i £ I I I | I I I I 250 300 Square Root Time [sec1/2] 5 2 0 - TISI J | f ■® I<5SI3 / ■ ------■r -i— [■ 1 - 1— i— i | i i i i | t ) * i i r 100 150 200 300 Square Root Time [sec1/2] Figure 71. Linear regression fit to a plot of thickness versus the square root of time for Ge-doped silicide coatings grown at 1150°C by packs comprised of (a) 12 wt.% si, 6% Ge, 2% MgF2, and sic, and (b) 16% Si, 8% Ge, 2% A1F3 and Al2o3. 177 Table 40. intrinsic parabolic rate constants (kp[cm2/s]) determined for the Ti(Si,Ge)2, Ti{Si,Ge), Ti5(Si,Ge)4, Ti5(Si,Ge}3, and Ti3(Si,Ge) layers grown by a pack comprised of 16 wt.% Si, 8% Ge, 2% A1P3 and A1203 on CP-titanium at 1150°C, 1050°C and 950°C, 1150°C 1050°C 950°C Ti(Si,Ge)2 1.31X10"10 3.39X10"11 4.58X10"12 Ti(Si,Ge) 5.61X10"11 1.16X10"11 7.57X10**13 Ti5(Si,Ge)4 5.96X10**11 3.04X10"11 8.93X10"13 Ti5(Si,Ge)3 5.54X10"12 2.26X10"12 7.57X10"13 Ti3 (Si,Ge) 5.96X10 -12 3.99X10”12 4.79X10"13 Table 41. Intrinsic parabolic rate constants (kp[cm2/s]) determined for the Ti(Si,Ge)2, Ti(Si,Ge), Ti5(Si,Ge}4, Ti5(Si,Ge}3, and Ti3(Si,Ge) layers grown by a pack comprised of 12 wt.% Si, 6% Ge, 2% MgF2 and Sic on CP-titanium at 1150°C, 1050°c and 950°C. 1150°C 1050°C 9 5 0 #C Ti(Si,Ge)2 1.57X10”10 9.22X10"11 1.42X10"11 Ti(Si,Ge) 4.24X10"11 2.17X10"11 4.83X10"12 Ti5(Si,Ge)4 6.39X10"11 2.70X10"11 1.31X10-11 Ti5(Si,Ge)3 6.41X10“12 3.30X10-12 1.06X10-12 Ti3 (Si,Ge) 7.56X10"12 3.45X10”12 1. 39X10”12 Table 42. Activation energies (Q[J/mol]) determined for the growth of the Ti(Si,Ge)2, Ti(Si,Ge), Ti5(Si,Ge)4, Ti5(Si,Ge)3, and Ti3(Si,Ge) layers by the KgF2- and AlF3-activated packs. HgF2-activated AlF3-activated Ti(Si,Ge)2 176,000 243,000 Ti(Si,Ge) 158,000 313,000 Ti5(Si,Ge)4 114,000 309,000 Tis(Si,Ge)3 131,000 144,000 Ti3(Si,Ge) 124,000 185,000 178 1150C 1050C 950C TiSi TiSi c *->to W c o CJ re DC ,-12 u Ti.Si ‘S5 e Ti Si *c c 0.00065 0.0007 0.00075 0.0008 0.00085 0.0009 <•> 1/T IK 1! 1150C 1050C 950C ~ i — ro E o TiSi. 4-> c CO to c TiSi o U 4->0) CO OC U 10’12-j *5) c *c -13 £ 10 0.00065 0.0007 0.00075 0.0008 0.00085 0,0009 (b) 1/T ricM Figure 72. A linear regression fit to an Arrhenius plot of log intrinsic rate constant versus inverse temperature for the Ti(Si,Ge)2, Ti(Si,Ge), Ti5(Si,Ge)4 , Ti5(Si,Ge)3, and Ti3(Si,Ge) layers grown by (a) MgF2-activated pack and (b) AlF3-activated pack. 179 volume for the germanide phase. The molar volumes for the Ti(Si,Ge) and Ti5(Si,Ge)4 solid solutions were estimated from the TigGeg phase. The intrinsic rate constants for the Ge- doped silicide phases grown by an A1F3- and MgF2-activated pack are reported in Tables 40 and 41, respectively. Arrhenius plots for the intrinsic rate constants of the Ge- doped silicide coatings grown by A1F3- and MgF2-activated packs are shown in Fig. 72. Theactivation energy for each of the silicide phases is reported in Table 42. The difference between the intrinsic parabolic rate constants for the Ti(Si,Ge)2, Ti5(Si,Ge)4, Ti5(Si,Ge)3, and Ti3(Si,Ge) layers grown by a MgF2-activated pack {Table 41) and undoped layers grown by a MgF2-activated pack (Table 14) is less than a factor of 2. However, the intrinsic rate constants for the growth of Ti(Si,Ge) are a factor of two slower than the undoped layer for a MgF2-activated pack. The activation energies for the growth of the Ti(Si,Ge)2, Ti5(Si,Ge)3, and Ti3(Si,Ge) layers by a MgF2-activated pack are approximately equal to the activation energies for the growth of TiSi2, Ti5Si3, and Ti3Si in the solid-state diffusion couple (Table 10). Since the intrinsic parabolic rate constants for the growth of TiSi2, Ti5Si3, and Ti3Si by the diffusion couple are similar to the growth rates for the respective Ti(Si,Ge)2, Ti5(Si,Ge)3, and Ti3(Si,Ge) layers by a MgF2-activated pack (Tables 9 and 41), solid-state diffusion of Si is the rate- limiting step. However, gas-phase diffusion was determined to 180 be the rate-limiting step for the growth of undoped and B- doped silicide coatings by a MgF2-activated pack, which produced low Si fluoride vapor pressures. The addition of Ge to the pack does not significantly change the gas-phase diffusion kinetics for the Si fluorides. Thus, the Ge additions affected the solid-state diffusion kinetics for the Ge-doped Ti-silicide phases, and changed the rate-limiting step for the growth of the Ti(Si,Ge)2, Ti5(Si,Ge)3, and Ti3(Si,Ge) layers from gas-phase diffusion control to the solid-state diffusion of Si. The Ge concentration profiles, which result from demixing in a chemical potential gradient, demonstrate that the diffusion rate of Ge is faster than Si. This observation does not imply that the interdiffusion rate of a Ge-doped Ti- silicide is faster than an undoped silicide. Tables 37 to 39 show that higher Ge additions produced thinner Ge-doped silicide coatings for the same coating cycle, which indicates that larger Ge-doping slowed the interdiffusion rate. A clear trend cannot be resolved by comparing the intrinsic rate constant for the Ge-doped silicide coating, undoped silicide coating and Ti-Si diffusion couple. The rate limiting step was defined by Wagner [71] as the slowest reaction step with the largest change in chemical potential. Therefore, the Ge- doping may have produced a large change in chemical potential, and comparing the magnitudes of the reaction rates is not sufficient to determine the change in mechanism. 181 The activation energies for the growth of the Ti(Si,Ge) and Tis(Si,Ge)4 layers are different than those determined for the undoped coating grown by a KgF2**activated pack and those of solid-state diffusion. However, the intrinsic parabolic rate constants for the growth of the Ti5(Si,Ge)4 phase are similar to the values measured for solid-state diffusion. The intrinsic parabolic rate constants for solid-state diffusion in TiSi are at least a factor of two faster than Ti(Si,Ge), which demonstrates that the Ge additions slow the growth of this layer. Thus, the growth of Ti(Si,Ge) and Ti5(Si,Ge)4 is probably controlled by solid-state diffusion, but the rate- limiting step is solid-state diffusion of Si that is slowed by the Ge solute. The ratio of the intrinsic parabolic rate constants for the growth of the Ge-doped silicide (Table 40) and undoped silicide (Table 16) coating layers for an AlF3-activated pack is a factor of 2 or 4. The activation energies for the growth of the Ti(Si,Ge )2 1 Ti(Si,Ge), and Ti5(Si,Ge)4 layers (Table 42) are significantly different than the undoped silicide layers (Table 17). The inclusion of the A1203 and A1F3 particles into the silicide coating probably tainted the kinetic measurements for the Ti(Si,Ge)2 layer. The differences in the kinetic data and morphologies between the Ti(Si,Ge)2 layer and TiSi2 layer grown by an A1F3 activator demonstrates that the growth mechanism has been changed. The rate-limiting step for the growth of the outer Ti(Si,Ge)2 182 layer could not be determined from the kinetic data. The intrinsic parabolic rate constants for the Ti(Si,Ge) and Ti5(Si,Ge)4 layers grown by A1F3- and MgF2-activated packs were similar at 1150°C and 1050°C, but a large difference was observed at 950°C. Thus, a change in the growth mechanism for the Ti(Si,Ge) and Ti5(Si,Ge)4 layers may occur at 950°C. Further work is required to resolve these differences. The intrinsic parabolic rate constants and activation energies for the Ti5(Si,Ge)3 and Ti3(Si,Ge) layers grown by an AlF3-activated pack are similar to the data for the undoped silicide coating and the solid-state diffusion couple. Thus, the rate-limiting step for the growth of the Ti5(Si,Ge)3 and Ti3(Si,Ge) layers is the solid-state diffusion of Si, in agreement with the results for the MgF2-activated pack. 4.5.4 Ge-Doped Silicide Coatings on Ti-Al-Nb Alloys The approach used to develop Ge-doped silicide coatings for the Ti-Al-Nb alloys was the same as that used for the B- doped silicide coatings, i.e. the results of the work on CP- titanium were applied to the Ti-Al-Nb alloys. Figure 73 shows the Ge-doped silicide coating grown on a Ti-22Al-27Nb alloy by a pack comprised of 12 wt.% si, 6% Ge, 2%MgF2 and sic at 1150°C for 12 hours. The coating consists of a thick outer layer of TiSi2 and TiSi, with a thin layer of a ternary Ti-Al- Nb phase adjacent to the substrate. The TiSi2 phase is the metastable C49 polymorph (JCPDS# 10-225, S.G. #63), with NbSi2 (JCPDS# 8-450) and minor TiAl3 peaks detected by XRD (Appendix Figure 73. SEM micrograph of the polished cross-section of a Ge-doped silicide coating grown on a Ti-22Al-27Nb alloy by a pack composed of 12 wt.% Si, 6% Ge, 2% M g F 2 and Sic at 1150°C for 12 hours. lOOum Figure 75. Optical micrograph of the polished cross-section of a Ge-doped silicide coating grown on a Ti-22Al-27Nb alloy by a pack comprised of 16 wt.% Si, 8% Ge, 2% A1F3 and A1203 at 1150°C for 12 hours. 184 100 7iSI2 TiSi H-22AI--27Nb * * * Base K E Silicon Profile o JSt 60 ^0 0—0-0-0 0 > o \ £n \ ___ o a Al Profi e E *AnA— o o Niobium Profile 10 ... .a. . .a. a. . *» *0* • • il 0 rA—t—A (a) 10 20 30 40 SO 60 70 80 90 100 110 120 130 140 ISO Distance From Outer Edge of Coating (jim) 100 90 .. TiSio < -TiSi Ti-22AI-27Nb Substrate S * 00 E 70 o • • Silicon Pilof le & 60 j>ocP0-a^ a 1 o 50 *3 S 40 a. Aluminum Profile E 30 -B o- o ’A'“ ■ "A* ° 20 hfl 10 Niobium Profile 0 ! Ai P 1Q-e»t<»-e»iCP_a^_Q-Q|-QQ 0P| O'l 0 50 100 150 200 250 300 350 400 450 500 550 (b) Distance From Outer Edge of Coating (/xm) Figure 74. EDS profiles for Ge-doped silicide coating grown on Ti-22Al-27Nb by a MgF,-activated pack at 1150°C for 12 hours, (a) Si, Ge, A1 and Nb profile and (b) Si, Al, and Nb profile. 185 A, Fig. 150a). The EDS profiles In Fig. 74 show that the Nb distribution within the coating was the same as that observed for the B-doped silicide coatings, which indicates the presence of NbSi2 precipitates. The Ge concentration in the Ti(Si,Ge)2 layer was highest at the Ti(Si,Ge)2/Ti(Si,Ge) interphase boundary, which correspond to the observations for the Ge-doped silicide coatings on CP-titanium. The Ge concentration is a little higher for the Tl(Si,Ge) layer, and no gradient for Ge was observed within the Ti(Si,Ge) layer. Figure 74b shows the wide interdiffusion zone with a 30 at.% enrichment of aluminum at the coating/alloy interface and depletion of Nb, in agreement with the interdiffusion zone observed for the B-doped silicide coating. Figure 75 shows the Ge-doped silicide coating grown on a Ti-22Al-27Nb alloy by a pack treatment composed of 16 wt.% Si, 8% Ge, 2% A1F3 and A1203. The coating includes a thick outer layer of Tisi2, which is the metastable C49 polymorph, with an inner layer of TiSi (Appendix A, Fig. 150b). The thickness of the inner Ti(Si,Ge) layer for the AlF3-activated pack is greater than the Ti(Si,Ge) layer observed for a MgF2-activated pack. A thick film of salt-oxide was observed at the surface of the Ti(Si,Ge)2 layer and some inclusions of activator and filler were observed in the surface region of the Ti(Si,Ge)2 coating, in agreement with the results observed for coatings grown on CP-titanium by the same pack. The A1F3 and Al2o3 inclusions at the surface region of Ti(Si,Ge)2 are indicated 186 100 1-22A1-27 Substrate Silicon Profile 1 a c .2 2 'to o Q. E <—Aluminum Profile o o Niobium Profile Ge Profil •jfi i.i ^ y 10 20 30 40 50 60 70 80 90 100 110 120 (a) Distance From Outer Edge of Coating (/im) too <■ -TiSi T1-22AI—27Nb Substrate « E o a Silicon Profile c o *.g *5 o CL Aluminum Profile E o o □ — □* ■H Niobium Profile 50 100 150 200 250 300 350 400 450 500 (b) Distance From Outer Edge of Coating G^m) Figure 76. EDS profiles of Ge-doped silicide coating grown on Ti-22Al-27Nb by an AlF3-activated pack at 1150°C for 12 hours, (a) Si, Ge, Al and Nb profile and (b) Si, Al, and Nb profile. 187 in Fig. 76a by the higher Al concentration at the surface of the Ti(Si,Ge)2 layer. Figure 76a shows that the amount: of Ge in the coating grown by the AlF3-activated pack is lower than that observed in the coatings grown by a MgF2 activator. This result was also observed for Ge-doped silicide coatings grown on CP- titanium, and is explained by the deposition of the salt- activator layer. The trends observed for the Ge concentration profiles within the Ti(Si,Ge)2 and Ti(Si,Ge) layers were the same as those shown for the MgF2-activated pack. The EDS profile of the substrate shows that the interdiffusion zone is not as wide, but a 30 at.% enrichment of Al was detected at the coating/metal interface. Ge-doped silicide coatings were grown on a Ti-20Al-22Nb alloy by pack treatments composed of 16 wt.% Si, 8% Ge, A1203 filler with either 2% A1F3 or 2% CuF2 at 1150°C for 12 hours. The features and morphologies for these coatings were the same as those observed for the Ge-doped silicide coating grown by the AlF3-activated pack on the Ti-22Al-27Nb alloy. The Ge concentrations for the coatings grown on the Ti-22Al-27Nb and Ti-20Al-22Nb alloy by the AlF3-activated pack were the same. The width and Al enrichment {30 at.% at the coating/alloy interface) of the interdiffusion zone was the same for coatings that were grown on either alloy substrate. The CuF2- activated pack produced larger amounts of Ge in the silicide coating, as predicted by thermodynamic calculations (Fig. 64). 188 The convolutions observed at the corners of the Ge-doped silicide coatings on either alloy substrate were similar to those produced by the B-doped silicide coatings. To remove the convolutions and reduce the width of the interdiffusion zone, a coating cycle of 950°C for either 6 or 28 hours was used for the three Ge-doped silicide coatings: (1) 12 wt.% Si, 6% Ge, 2% HgF2 and SiC, (2) 16 % Si, 8% Ge, 2% A1F3 and A1203, and (3) 16 % Si, 8% Ge, 2% CuF2 and A1203. The thin, adherent coating produced by an A1F3- or CuF2-activated pack at 950°c for 6 hours on Ti-20Al-22Nb is shown in Figs. 77a and 77b, respectively. Note that a salt layer is still observed at the surface of the Ti(Si,Ge)2 layer, and a thin Ti(Si,Ge) layer was detected adjacent to the coating/alloy interface. The thicknesses for the Ge-doped silicide coatings grown on Ti- 22Al-27Nb and Ti-20Al-22Nb at 1150°C or 950°C by the MgF2-, A1F3- and CuF2-activated packs are summarized in Tables 43 to 45. The EDS profiles for the coatings grown by MgF2-, A1F3- and CuF2-activated packs on Ti-20Al-22Nb are shown in Figs. 78 to 80. The width of the interdiffusion zone is smaller for each Ge-doped silicide coating, and an Al enrichment of 50 at.% is observed at the coating/alloy interface, in agreement with the B-doped coatings. The highest Ge concentration within the Ti(Si,Ge)2 layer of the thinner coatings was observed at the Ti(Si,Ge)2/Ti(Si,Ge) interface. The Ge concentration in the thinner coatings grown by the A1F3- and 189 Ba>«*Altoy Figure 77. SEM micrograph of the polished/etched cross- section of a Ge-doped silicide coating grown on a Ti-20Al-22Nb alloy at 950°C for 6 hours by a pack comprised of (a) 16 wt.% Si, 8% Ge, 2% A1F3 and A1203, and (b) 16 wt.% Si, 8% Ge, 2% Cu F2 and A1203. 190 Table 43. The average thicknesses ([/xm]> for the Ti(Si,Ge)2 and Ti(Si,Ge) layers grown on Ti-22Al-27Nb and Ti- 20Al-22Nb by a pack comprised of 12 wt.% Si, 6% Ge, 2% MgF2, and Sic at three coating treatments: (1) 1150°C for 12 hours on Ti-22Al-27Nb (2) 950#C for 28 hours on Tl-20Al-22Nb (3) 950°C for 6 hours on Tl-20Al-22Nb 1150°C 950#C 950°C 12 hours 28 hours 6 hours Ti{Sl,Ge)2 123.80 30.37 17.75 Tl(Si,Ge) 19.43 5.41 4.04 100 TiSi.TiSi Ti-20Al-22Nb Base Alloy 80-i Si Profile £ ,ca. a 60 ^ o Al Profile 'U> a 40*; o6 Ge Profile U __ 2 0 -= O*’ Nb Profile 0 10 20 40 5030 Distance [pro] Figure 78. EDS Si, Ge, Al and Nb profiles for a Ge-doped silicide coating grown on Ti-20Al-22Nb by a MgF2-activated pack at 950°c for 6 hours. 191 Table 44. The average thicknesses ([/Ltm]) for the Ti(SirGe)2 and Ti(Si,Ge) layers grown on Ti-22Al-27Nb and Ti- 20Al-22Nb by a pack comprised of 16 wt.% Slr 8% Ge, 2% A1F3, and A1203 at three coating treatments: (1) 1150°C for 12 hours on Ti-22Al-27Nb (1) 1150°C for 12 hours on Ti-20Al-22Nb (2) 950 °C for 28 hours on Ti-20Al-22Nb (3) 950°C for 6 hours on Ti-20Al-22Nb 1150°C 950°C 950°C 12 hours 28 hours 6 hours Ti-22Al-27Nb Tl(Sl,Ge)2 48.90 -- -- Ti(Sl,Ge) 29.85 -- -- Tl-20Al-22Nb Tl(Sl,Ge)2 39.72 27.39 16.04 Ti(Sl,Ge} 26.33 6.42 4.28 100 TiSi, TiSi Ti-20Al-22Nb j Base Alloy • 80 -E Si Profile 60-; Al Profile *co 8. 40-E Ge Profile 2 0 -: Nb Profile 10 15 20 25 30 35 Distance [fim] Figure 79. EDS si, Ge, Al and Nb profiles for a Ge-doped silicide coating grown on Ti-20Al-22Nb by an AlF3-activated pack at 950°c for 6 hours. 192 Table 45. The average thicknesses ([frn]) for the Ti(Si,Ge) and Ti(Si,Ge) layers grown on Ti-22Al-27Nb and Ti I KJ 20Al-22Nb by a pack comprised of 16 wt.% Si, 8% Ge, 2% CuF2, and A1203 at three coating treatments: (1) 1150°C for 12 hours on Ti-20Al-22Nb (2) 950®C for 28 hours on Ti-20Al-22Nb (3) 950°C for 6 hours on Ti-20Al-22Nb 1150°C 950°C 950°C 12 hours 28 hours 6 hours Ti(Si,Ge)j 48.93 22.04 7.60 Ti(Si,Ge) 29.18 7.81 3.53 100 TiSi, TiSi Ti-20Al-22Nb Base Alloy 80 H Si Profile 60-i Al Profile C/1 a Ge Profile Nb Profile Distance Qim] Figure 80. EDS Si, Ge, Al and Nb profiles for a Ge-doped silicide coating grown on Ti-20Al-22Nb by a CuF,-activated pack at 950°c for 6 hours. 193 CuF2-packs was similar to that observed for the thicker coatings. The Ge concentration was highest for coatings grown by a CuF2-activated pack and lowest for the coating grown by a MgF2-activated pack, in agreement with the predictions based on thermodynamic calculations. However, the Ge-doping for the thick coatings formed at 1150°c for 12 hours was highest for the MgF2-activated packs. The salt-oxide layers at the surface of the thinner coatings grown by the A1F3- and CuF2- activated packs were thin, and inclusions at the surface of Ti(Si,Ge)2 were not observed. Therefore, the reason that thick Ge-doped silicide coatings grown on CP-titanium and Ti- Al-Nb alloys by a MgF2-activated pack contain more Ge than packs grown using A1F3 and CuF2 activators, which contradicts the thermodynamic calculations, is the A1203 filler for these packs. The activator-oxide deposits that collect at the surface of the silicide coating limit the influx of Ge into the coating, which results in lower Ge concentrations when an A1203 filler is used in the pack. 4.6 Ge-Dopsd and Undoped Molybdenum-Bilicide coatings The development, growth and high-temperature oxidation resistance for these coatings have been described by Mueller et al. [25-27]. The more detailed study in this work yielded new results that have not yet been published. These issues are addressed in the following section. 194 4.6.1 Characterisation of tha Coatings A cross-section of the Ge-doped Mo-silicide coating grown by a pack composed of 10 wt.% Si - 10% Gef 2% NaF and Sic (wt%) is shown in Fig. 81. The thick outer layer of the coating in Fig. 81a is Mo(Si,Ge)2, and EDS revealed that the thin layers adjacent to the molybdenum substrate in Fig. 81b are Mo5(Si,Ge)3 and Mo3(Si,Ge). [27] The EDS profile in Fig. 82 indicates that 1-9 at.% Ge is distributed in the coating; the Ge concentration in Mo(Si,Ge)2 is highest at the Mo5(Si,Ge)3/Mo(Si,Ge)2 interphase boundary, and decreases in the Mos(Si,Ge)3 and Mo3(Si,Ge) layers. The chemical demixing of the Mo(Si,Ge)2 solid solution by the chemical potential gradient establishes this composition profile for Ge [25]. Schmalzried and Laqua [103] showed that demixing of a multicomponent (A,B)0 oxide occurs in a oxygen potential gradient, with an enrichment of AO at the side of the highest oxygen potential if DA>DB» D 0. Similarly, the observed chemical demixing for the Mo(Si,Ge)2 solid solution in the Mo (or Si) chemical potential gradient would result if Doe>Dsi>>DMo' giving Ge enrichment at the Mo(Si,Ge)2/Mos(Si,Ge)3 interphase boundary. The reported parabolic rate constants for the growth of a single-layered MoGe2 diffusion coating on Mo [155] are 10 to 15 times higher than the single-layered growth rate for MoSi2 reported by Gage and Bartlett [94] at 950 to 1150°C. Therefore, the solid- state diffusion rate for Ge in Mo(Si,Ge)2 is much faster than 195 Si, which indicates that the observed profile for Ge in Mo(Si,Ge)2 is the result of chemical demixing. Figure 83 shows a SEM micrograph of the surface of a Ho(Si,Ge)2 coating grown by a NaF-activated pack. Cracks are visible at the surface of the coating, and the small surface protrusions marked "a" in Fig. 83 are sodium-rich oxide deposits. These deposits were extremely adherent, and could not be removed by vigorous cleaning in water, acetone, or ethanol. Similar byproduct layers were observed on MoSi2 coatings grown with other activators, and the composition of the layer reflected the specific halide activator used to grow the coating. The cross-sections of superficial salt layers on the MoSi2 coatings grown by NaF- and MgF2-activated packs are shown in Figs. 84a and 84b. The approximate elemental compositions of these layers were determined by EDS and are summarized in Table 46. A Na-rich oxide deposit results from using a NaF activator, and a Mg-rich oxide deposit results from MgF2. The K and Ca impurities were present in the sic powder used as the pack filler. The A1 impurity results from the fluoride vapors reacting with the A1203 crucible to form A1F3 vapor, which deposit A1 at the surface. The same type of impurity salt layer was observed on coatings formed with and without the addition of Ge. Byproduct layers present on the surfaces of coatings grown by A1F3, CuF2, and MnF2 activators also reflected the composition of the activator, i.e., an Al- rich salt oxide layer was observed on coatings grown by an 196 Ho(Si,Ge)2 Figure 81. Cross-sections of a Ge-doped MoSi2 coating grown by a pack composed of lOSi-lOGe, 2NaF and Sic (weight%) at 1150°C for 12 hours, (a) SEM micrograph showing Mo(Si,Ge)2 and Ho 5 (Si,Ge)3 , and (b) SEN micrograph showing the thin layers of M o 5 (Si,Ge)3 and Mo3 (Si,Ge)* 197 K o E Mo(Si,Ge) o 8-. Ge Profile aE 7 - c □ 6 ■ ■ E v o o c o crj O CL OE o I 20 30 40 SO 60 70 80 Distance From Outer Edge of Coating (pm) Figure 82. EDS profile of germanium in the three-layered Mo(si,Ge)2/Mos(Si,Ge)3/Mo3(SifGe) coating grown by a NaF- activated pack at 1150°C for 12 hours. Figure 83. SEM micrograph of the surface of a Ge-doped silicide coating grown by a NaF activated pack. The protrusions marked "a" are sodium rich. 198 Table 46. Elements present in the byproduct layers on the surface of MoSl2 coatings grown by NaF- and MgF2- activated packs (Pack# 1 and 3). Major: 15-60 at.%, Medium: 3-10 at.%, Minor 0.5-2 at.%. NaF-pack MgF2-pack Major Al, o, si Al, 0, si Medium F, Ha F, Mg Minor Ca, K, Mo Ca, K, Mo Figure 84. High-magnification SEM micrograph of the byproduct salt-oxide layers at the surface of MoSi2, (a) coating grown by a NaF-activated pack and (b) a MgF2-activated pack. 199 AlF3-activated pack, etc. Byproduct salt layers were previously observed by Levine and Caves [66] and Kandasaray et al. [85] for the aluminizing of nickel-base superalloys by condensed halide activators. Since all the activators used in this study are in a condensed state at 950 to 1150°C, the deposition of byproduct layers is not surprising. Computer-assisted, thermodynamic calculations for the equilibrium vapor pressures that result from a pack composed of pure Si, pure Ge, NaF and Sic (see Table 47) demonstrate that the partial pressures of Na and NaF vapors are high enough to expect condensation at the coating surface. Page and Bartlett [94] studied the growth kinetics for MoSi2 diffusion coatings grown by pack cementation on pure Ho, and used marker experiments to show that the inward diffusion of silicon was dominant. These authors did not mention the presence of salt-oxide deposits at the surface. Therefore, due to the inward growth of these coatings, impurities were not observed in the bulk of the molybdenum-silicide layers, but accumulated at the surface as a convenient marker. 4.6.2 Growth Kinetics of Go-Doped Molybdenum Bilioide The parabolic rate constant for the diffusion-controlled growth of single-layered undoped MoSi2 coatings was determined by Page and Bartlett [94]. However, the coating shown in Fig. 81 consisted of the three layers Mo(Si,Ge)2, Mo5(Si,Ge)3, and Ho3(Si,Ge). Figures 85a, 85b and 85c present linear regression fits to the respective plots of layer thickness 200 Table 47. The calculated equilibrium partial pressures for pack comprised of pure Si, pure Ge, NaF activator and Sic filler at 1423 K. Partial Partial Pressure Pressure Ar 9.665X10-1 C 3.976X10**21 Na 2.388X10"2 c f 2 8.852X10”22 SiF4 5.161X10”3 C2 6.564X10"26 NaF 3.021X10 C3 2.085X10“26 SiF3 9.882X10"4 F2 9.218X10"28 Na2F2 3.227X10”4 cf3 9.497X10"29 SiF2 1.241X10"4 CF4 9.521X10"31 GeF2 2.961X10”® C2F4 2.260X10”31 Na2 1.604X10“® C4 1.281X10"34 Ge F 6.787X10"7 c5 9.746X10'37 Ge 2.111X10-7 C2F4 3.372X10-41 SiF 3.135X10“® C2F6 2.808X10-52 Si 1.437X10”® Si2C 1.247X10"12 1.090X10”12 si3 1.677X10-13 F 5.961X10”14 GeF3 6.696X10"16 SiC2 3.689X10-17 GeF4 2.449X1O"10 sic 2.100X10’19 CF 2.386X10”20 201 versus the square-root of tine for Mo(Si,Ge)2, Mo5(Si,Ge)3 and Mo3(Si,Ge) grown by a NaF-activated pack. The thickness of each layer was corrected by accounting for an incubation tine of each layer by Eq. (4.14). The correlation coefficient of each plot was at least 0.95, which indicates that the kinetics are parabolic. With the reasonable assunption of solid-state diffusion control, the theory of Hsu [29] and Wang et al. [30] (Eq. 2.26) was applied to deternine the intrinsic rate constants (kp) fron the neasured apparent rate constants (kp') for the triple-layer growth of Mo(Si,Ge)2, Mo5(Si,Ge)3 and Mo3(Si,Ge), 1 p 3 (4.22a) (4.22b) m o w * (4.22c) where 1 is Mo3Si, 2 is Mo5Si3, 3 is MoSi2, is the average molar volume, is the stoichiometry normalized to one mole of silicon, and is the average thickness of each layer. The molar volumes for the Ge-doped silicides were estimated by a rule-of-mixtures relationship, based on the average mole fraction of MovGe (XMovGe) in each layer and the assumption 202 that the averaged molar volume does not vary with the thickness: V[Mo(Si,Ge)2] = (XMoGe2)VMoCo2 (4.23a) V[Mo5(Si,Ge)3] ** (l“XMo50a3)VMo5si3 + (XM o50e3)VMo50o3 ( 4 .2 3 b ) V[Mo3(Si,Ge) ] - (l-XMo3Co)VMo3si + tXMo30o)VMo30o (4.23c) These molar volumes were used to determine the intrinsic rate constants for the Ge-doped silicide coatings using Eqs. (4.22a)-(4.22c) at 950, 1050 and 1150°C, as summarized in Table 48. For solid-state diffusion control and guasi-steady-state growth for the molybdenum-silicides driven by the chemical potential gradient within each layer, a modified form of the Wagner scaling theory (Eg. 2.16) should predict the intrinsic parabolic rate constants [110,111]. With the assumption that the averaged interdiffusion coefficient reported in the literature is independent of composition for each compound, Eg. (2.16) is extended to the three-layered case: Xp(MOi/2si) ■ D(MOjy2Si) {-lnag^jjj} (4.24a) Xp(Mosy3Si) “ D(Mo5^3Si) {lnasi|i]-lna£ij2] J1 (4.24b) kp(Mo3Si) = D(Mo3Si) {lnaslt2j-lnasl[3j> (4.24c) where is the interdiffusion coefficient for each phase, asi[l] activity at the MoSi2/Mo5Si3 interface, asij2j is the Si activity at the Mo5Si3/Mo3Si interface and asi[3] is 203 the Si activity at the Mo3Si/Mo interface. If solid-state diffusion is rate-controlling, then the intrinsic parabolic rate constant predicted by Eqs. (4.24a-4.24c) should be equivalent to the experimentally determined intrinsic rate constants of Eqs. (4.22a-4.22c). The atomic radius of the Ge atom is larger than that for Si, and the addition of Ge solute would be expected to change the diffusional growth kinetics of the silicide layers. To determine the effect of Ge, the intrinsic rate constants for a three-layered coating of undoped silicide were calculated using Eqs. (4.24a-4.24c). The interdiffusion coefficients for HoSi2 were calculated from the reported growth kinetics for a single-layered MoSi2 coating [94,118,119] using Eq. (2.16). The interdiffusion coefficients for Mo5Si3 and Mo3Si were determined from the diffusion data of Bartlett et al. [115] and Fitzer and Matthias [116], and the calculated intrinsic rate constants at 950, 1050 and 1150°C are given in Table 49. The thermodynamic data of Pankratz et al.[l5l] were used to determine the silicon activities required for the calculations by Eqs. (4.22a-4.22c). Steinmetz et al. [155] reported a parabolic rate constant for the growth of a single-layered MoGe2 coating. Isothermal interdiffusion between the MoGe2-Mo couple of this coating resulted in the formation of Mo13Ge23, Mo5Ge3 and Mo3Ge layers; apparent parabolic rate constants were reported at two temperatures. The intrinsic parabolic rate constants for the growth of these compounds were 204 determined using the form of Eqs. (4.20a-4.20c) and extrapolated to the temperatures of interest, as reported in Table 50. The Arrhenius plot in Fig. 86a compares the intrinsic parabolic rate constants for Mo(Si,Ge)2 with the MoSi2 rate constants calculated from the reported interdiffusion data and the reported rate constant for MoGe2. The activation energies for the growth of Mo(Si,Ge)2, MoGe2 and MoSi2 layers in a three-layered coating are given in Table 51. The intrinsic parabolic growth rate constant for Mo(Si,Ge)2 is a factor of 10 to 15 slower than that for MoGe2, but a factor of two faster than the rate constants for undoped MoSi2 reported by Page and Bartlett [94] and Cox and Brown [118]. The interdiffusion data for MoSi2 measured by Samsanov and Koval'chenko [119] are significantly different than the other data for MoSi2, and are not used for comparison. Since the activation energies for the growth of Mo(Si,Ge)2 and pure MoSi2 are the same, the mechanism for solid-state diffusion is not changed by the addition of Ge to MoSi2. The growth rate for MoGe2 is much faster than MoSi2 [94,118,155], and one may expect that alloying MoSi2 with a fast diffusing component would increase the interdiffusion coefficient. The germanium profile that results from chemical demixing during Mo(Si,Ge)2 growth indicates that DGB>Dsi» D Mo, which corresponds with expectation. The addition of the faster diffusing germanium to the MoSi2 compound explains the 205 Uo(SI.Ce). 1423K 1323K I223K Square Root Time (eec)1/ 2 3.0 1423K 1.0 1323K 1223K 0.0 0 SO 100 ISO 200 300 Square Root Time ( e e e ) ' / 2 0.3 0.4'- 1423K 0.3 0.2 ■ 1223K 0.1 0.0 SO 100 ISO 200 2S0 300 Square Root Time (eac)1/2 Figure 85. Linear regression fit to a plot of layer thickness versus the square-root of time for Ge-doped coatings grown by NaF-activated pack at 1150, 1050 and 950°C for (a) Mo(Si,Ge)2, (b) Mo5(Si,Ge)3, and (c) Mo3(Si,Ge). 206 Table 48. Intrinsic parabolic rate constants (kp [cm2/s]) determined from the experimental data for a Ge- doped silicide coating produced using a lOSi- lOGe/NaF/SiC pack at 950, 1050 and 1150°C. 1150°C 1050°C 950°C Mo(Si,Ge)9 4.65X10"10 1.41X10"10 3.03X10"13 MOe(Si,Geys 4.97X10”12 1.62X10"12 2.48X10-13 Mo3 (Si,Ge) 2.73X10"12 5.60X10"13 2.04X10**13 Table 49. Theoretical intrinsic parabolic rate constants (kp tcm2/s]) for the growth of a three-layer raolybdenum- silicide coating at 950, 1050 and 1150°C that are calculated from diffusion data reported in the literature. 1150°C 1050°C 950°C MoSi2 [94] 2.72X10 7.26X10"11 1.62X10"11 MoSi2 [118] 1.92X10"10 4.58X10"11 1.33X10-11 MoSi2 [119] 1.15X10“® 5.14X10-10 1.68X10"10 Ho5Sl3 [115] 5.26X10"11 5.78X10“12 4.39X10“13 HOeSil [116] 3.38X10"10 4.61X10"11 4.48X10"12 Ho3Si [115] 1.57X10"13 2.01X10-14 1.86X10-15 Table 50. Intrinsic parabolic rate constants (kp [cm2/s]) for the growth of a three-layer Mo-Ge coating at 950, 1050 and 1150°C that are calculated from the reported rate constants for a Mo-MoGe, diffusion couple [155], and the single-layer growth rate for MoGe2 reported in literature [155]. 1150°C 1050°C 950°C MoGe2 4.94X10“® 2.22X10“® 8.70X10”10 Mo13Si23 1.55X10“® 1.54X10“® 1.04X10-10 Ho5Ge 5 1.67X10"10 3.87X10”11 7.06X10"12 Mo3Ge 3.41X10-11 9.36X10"12 2.08X10-12 207 1150C 1050C 950C 10‘7 Y Y MoSi (Samsanov & Koval'chenko]enko] Z 10‘N MoGe: [Steinm etz e t al,} 2 1q.9 ■ w 10 ~- c o u o 101 A-10 M Mo(Si,Ge). oceg Mo SI2 [Page & Bartlett] f ^ •a«/) 10*"-J c ■c MoSI2 [Cox & Brown] •E 10* Ca) 0.00065 0.0007 0.00075 0.0008 0.000B5 0.0009 1/T TK*1! 1150C 1050C 950C 7 ----- ? £ E 10*91 Mo Ge [Steinmetz et al.] u \ Q [Fitzer & Matthias] Mo3Ge [Steinmetz et at. £c 10*,(\ n 0b>. ^ MOjSi j [Bartlett I/) -n ■ a 10 " o u 10*,z 4-» Figure 86. Arrhenius plot of the intrinsic parabolic rate constant for (a) the experimentally measured Mo(Si,Ge)2 outer layers compared with the calculated intrinsic rate constants for MoSi2 [94,118,119], and (b) measured intrinsic rate constants for the inner Mo5(Si,Ge)3 and Mo3(Si,Ge) layers compared with calculated rates for the Mo5Si3, Mo3Si, Mo5Ge3 and Mo3Ge layers [115,116,155]. 208 Table 51. Activation energies (Q [j/inol]) determined from the growth of each layer for the three-layered Ge-doped silicide coating, and calculated values for the rate constants for MoSi2, and Mo5Si3 and Mo3Si. Mo-Si-Ge Q[J/mol] Mo-Si Q[J/mol] Mo(Si,Ge), 198.000 MoSi2 [94] 204.000 Mo5(Si,Gef3 218.000 MoSi2 [118] 192.000 Mo3 (Si,Ge) 186,000 MoSi2 [119] 139.000 Mo5Si3 [115] 346.000 Mo5Si3 [116] 313.000 Mo3Si [115] 321.000 Mo-Ge Q[J/mol] * MoGe2 [155] 125.600 HosGe3 [155] 229,000 Mo3Ge [155] 202.600 faster intrinsic rate constant for Mo(Si,Ge)2 [104,105]. Furthermore, the melting point of pure MoSi2 (2293 K) is much higher than MoGe2 (1353 K) [57], and the addition of a solute with a lower melting point generally increases the rate of solid-state diffusion [104,105]. The Arrhenius plot in Fig. 86b compares the measured intrinsic rate constants for Mo5(Si,Ge)3 and Mo3(Si,Ge) with the rate constants calculated from the interdiffusion data for Mo5Si3, Mo3Si, Mo5Ge3 and Mo3Ge. The growth rate for Mo5(Si,Ge)3 is slower than Mo5Si3 and MosGe3, but the activation energy for Mo5(Si,Ge)3 growth given in Table 51 is approximately the same as for Mo5Ge3. The growth rate for Mo3(Si,Ge) is slower than Mo3Ge, but much faster than Mo3si; the activation energy for Mo3(Si,Ge) growth is almost the same as for Mo3Ge. similar to MoGe2 and MoSi2, the parabolic 209 growth rate constants for Mo5Ge3 and Mo3Ge are 5 to 20 tines faster than for Mo5Si3 and Mo3Si, respectively, and Ge would be expected to have a fast rate of diffusion. Chenical demixing that would result in the highest Ge concentration at the interface with the largest molybdenum chemical potential is not apparent in Fig. 82. However, these inner compound layers were too thin to accurately resolve the Ge profiles within each layer. The crystal structure of Mo5Si3 (tetragonal; tl32, S.G. #140) is the same as MosGe3, and Mo3si (cubic; cP8, S.G. #223) is the same as Mo3Ge [27,57]. Therefore, the Si and Ge atoms of Mo5(Si,Ge)3 and Mo3(Si,Ge) solid solutions certainly diffuse on the same sublattice, and simple chemical demixing would be expected. THE DEVELOPMENT, GROWTH AND OXIDATION RESISTANCE OP BORON- AND GERMANIUM-DOPED SILICIDE DIFFUSION COATINGS BY FLUORIDE-ACTIVATED PACK CEMENTATION VOLUME II DISSERTATION Presented in Partial Fulfillment of the Requirements for the Degree Doctor of Philosophy in the Graduate School of The Ohio State University by Brian Vern cookeram, B.s., M.S. The Ohio State University 1994 Dissertation Committee: Approved by R.A. Rapp K.H. Sandhage &Udta. Advisor W. Soboyejo Program in Department of Materials Science and Engineering TABLE OF CONTENTS CHAPTER PAGE V. OXIDATION RESISTANCE ON BORON- AND GERMANIUM- DOPED SILICIDE COATINGS: RESULTS AND DISCUSSION...... 210 5.1 Isothermal oxidation ...... 210 5.1*1 Boron-Doped Silicide Coatings on CP-Titanium ...... 210 5.1.2 Uncoated Ti-22Al-27Nb ...... 229 5.1.3 Boron-Doped Silicide Coating on Ti-22Al-27Nb ...... 233 5.1.4 Germanium-Doped Silicide Coating on CP-Titanium ...... 239 5.1.5 Germanium-Doped Silicide Coating on Tl-22Al-27Nb ...... 247 5.1.6 Low Temperature oxidation ...... 252 5.2 Prevention of Posting for MoSi2 coatings and Bulk MoSi2 ...... 252 5.2.1 Undoped and Ge-doped MoSi2 Diffusion Coatings...... 253 5.2.2 Posting Resistanceof Bulk MoSi2 . . . 263 5.3 Cyclic Oxidation ...... 270 5.3.1 Boron-Doped Silicide Coatings on CP-Titanium ...... 270 5.3.2 Boron-Doped Silicide Coating on Ti-22Al-27Nb ...... 288 5.3.3 Germanium-Doped Silicide Coating on CP-Titanium ...... 302 5.3.4 Germanium-Doped Silicide Coating on Ti-22Al-27Nb ...... 318 VI. SUMMARY AND CONCLUSIONS ...... 333 VII. FUTURE WORK ...... 335 xxviii APPENDICES A. X-RAY DIFFRACTION OF AS-COATED WORKPIECES. . . . 336 B. THICKNESS RATIOS FOR THE GROWTH OF UNDOPED SILICIDE COATINGS ...... 343 C. PARABOLIC FIT TO THE ISOTHERMAL OXIDATION KINETICS FOR THE BORON- AND GERMANIUM-DOPED SILICIDE COATINGS ...... 345 D. X-RAY DIFFRACTION FOR BORON- AND GERMANIUM-DOPED SILICIDE DIFFUSION COATINGS FOLLOWING ISOTHERMAL OXIDATION ...... 355 E. CYCLIC OXIDATION DATA ...... 363 REFERENCES ...... 405 xxix LIST OF TABLES TABLE PAGE 52. The correlation coefficients (R2) determined from the linear regression fit for a parabolic (dw vs t1'2) and log(dw)-log(t) plot (with the power term (n)) for B-doped silicide coating formed on CP-titanium by packs composed of (a) 7 wt.% Si, 6% TiB2, 2% MgF2 and A1203, and (b) 7 wt.% Si, 6% TiB2, 2% MgP2 and A1203...... 217 53. Activation energies (Q[J/mol]) determined for the isothermal oxidation of the B-doped silicide coatings grown on CP-titanium and Ti-22Al-27Nb at 1150°C by packs composed of (a) 7 wt.% Si, 6% TiB2, 2% MgF2, A1203, (b) 7% Si, 6% TaB2, 2% MgF2, Al203 at 1150°C for 12 hours, and reported data for "the oxidation of Si, Tisi2, Ti, Nb and TiB2 [10,41,56,156]. The correlation coefficients (R2) for the fit to the Arrhenius plot are also reported with the proposed mechanism for oxidation...... 218 54. Activation energies (Q[J/mol]) determined for the isothermal oxidation of uncoated Ti-22Al-27Nb and reported data for the oxidation of TiSi2, Nb, Ti, a2-base Ti-Al-Nb alloys and A1203 forming alloys [10,56,161] ...... 235 55. Activation energies (Q[J/mol]) determined for the isothermal oxidation of the Ge-doped silicide coatings grown on CP-titanium and Ti-22Al-27Nb by a pack comprised of either; (a) 12 wt.% si, 6% Ge, 2% MgF2, SiC at 1150°C for 12 hours, or (b) 16% Si, 8% Ge, 2% A1F3, A1203 at 1150°C for 12 hours and reported data for the oxidation of Si, TiSi2, Nb, Ti and TiB2 [10,41,56,156]. The correlation coefficients (R2) for the Arrhenius fit are given with the proposed mechanism of oxidation...... 244 56. The weight change observed for B-doped and Ge-doped silicide coatings after 2000 hours of isothermal oxidation at 500°C...... 255 XXX 57. The weight change that results from application of the salt layer from an aqueous solution or Si(OC2H5)4, Borax (Ha2B407) or sodium silicate solution...... * . . . . 266 58. The weight gain measured following 500 hours of isothermal oxidation at 500"C for a MoSi2 heating element coated with NaN03, Si(OC2He)4, Borax (Na2B407) or sodium silicate solution ...... 272 59. A summary for the detection of oxygen penetration by microhardness and etching for B-doped silicide coatings grown on CP-titanium by six different packs at 1150°C for 12 hours following cyclic oxidation at 500°C, 600#C, 700°C, 800°C, 900#C and 1000°C. The number in brackets is the number of coupons tested: (1) Si-TiB2/MgF2/Al203 (2) Si-TaB2/MgF2/Al203 (3) Si-TiB2/AlF3/Al203 (4) Si-TaB2/AlF3/Al203 (5) Si-TiB2/CuF2/Al203 (6) Si-TaB2/CuF2/Al203 . . . .278 60. A summary for the detection of oxygen penetration by microhardness and etching for the undoped silicide coatings grown on CP-titanium by four different packs at 1150°C for 12 hours following cyclic oxidation at 500°C, 600°C, 700°C, 800°C, 900°C and 1000°C. Two coupons were tested at each temperature: (1) Si/MgF2/Al203, (2) Si/A1F3/Al203, (3) Si/CuF2/Al203, and (4) Si/MgF2/SiC ...... 285 61. A summary for the detection of oxygen penetration by microhardness and etching for the thin B-doped silicide coatings grown on CP-titanium by two different packs at either 950°c for 6 hours or 950°C for 28 hours following cyclic oxidation at 500°C, 600°C, 700°C, 800°C, 900°C and 1000°C, with the number of coupons tested at each temperature in brtickofcs• (1) Si-TiB2/MgF2/Al203 at 950°C for 6 hours (2) Si-TiB2/CuF2/Al203 at 950°C for 6 hours (3) Si-TiB2/MgF2/Al203 at 950°C for 28 hours (4) Si-TiB2/CuF2/Al203 at 950°C for 28 hours . . . .287 xxxi 62. A summary for the detection of oxygen penetration by microhardness and etching for the B-doped silicide coatings grown on Ti-22Al-27Nb (22-27) and Ti-20Al-22Nb (20-22) by the pack treatments below following cyclic oxidation at 500-1000°C, with the number of coupons tested at each temperature in brackets: (1) Si-TiB2 /MgF2/Al20 3 on Ti-22Al-27Nb at 1150°C, 12h (2) Si-TiB2/MgF2/Al203 on Ti-20Al-22Nb at 1150°C, 12h (3) Si-TiB2/MgF2/Al203 on Ti-20Al-22Nb at 950°C, 28h (4) Si-TiB2/MgF2/Al203 on Ti-22Al-27Nb at 950°C, 6h (5) Si-TiB2/MgF2/Al203 on Ti-20Al-22Nb at 950°C, 6h (6) Si-TaB2/MgF2/Al203 on Ti-20Al-22Nb at 1150°C, 12h (7) Si-TiB2/CUF2/Al203 on Ti-20Al-22Nb at 950°C, 28h (8) Si-TiB2/CuF2/Al203 on Ti-20Al-22Nb at 950#C, 6h (9) Si-TaB2/CuF2/Al203 on Ti-20Al-22Nb at 950#C, 28h (10) Si-TaB2/CuF2/Al203 on Ti-20Al-22Nb at 950°C,6h...... 295 63. A summary for the detection of oxygen penetration by microhardness and etching for Ge-doped silicide coatings grown on CP-titanium by three different packs at 1150°C for 12 hours following cyclic oxidation at 500°C, 600°C, 700®C, 800°C, 900°C and 1000°c. The number in brackets is the number of coupons tested: (1) Si-Ge/MgF2/SiC, (2) Si-Ge/A1F3/Al203, (3) Si-Ge/CuF2/Al203 ...... 304 64. A summary for the detection of oxygen penetration by microhardness and etching for Ge-doped silicide coatings grown on CP-titanium by three different packs with two ratios of Si to Ge at 1150°C for 12 hours following cyclic oxidation at 500®C, 600#C, 700°C, 800°C, 900°C and 1000°C. Two coupons were tested for each pack: (1) 12 wt.% Si, 12% Ge, 2% MgF2, SiC[l:l] (2) 6% si, 12% Ge, 2% MgF2, Sic [1:2] (3) 16% Si, 16% Ge, 2% A1F3, A1203 [1:1] (4) 8% Si, 16% Ge, 2% A1F3, A1203 [1:2] (5) 16% Si, 16% Ge, 2% CuF2, A1203 [1:1] (6) 8% Si, 16% Ge, 2% CuF2, A1203.. [1:2]...... 311 xxxii 65. A summary for the detection of oxygen penetration by tnicrohardness and etching for Ge-doped silicide coatings grown on CP-titanium at 950°C for either 28 hours, 12 hours or 6 hours by three different packs following cyclic oxidation at 500°C, 600°c, 700°C, 800°C, 900°c and 1000°C. Two coupons were tested for each pack: (1) Si-Ge/MgF2/SiC, 950°C, 28 hours (2) Si-Ge/MgP2/SiC, 950°C, 6 hours (3) Si-Ge/A1F3/Al203 950aC, 28 hours (4) Si-Ge/A1F3/Al203 950aC, 12 hours (5) Si-Ge/CuF2/Al203 950°C, 28 hours (6) Si-Ge/CUF2/Al203 950°C, 12 h o u r s ...... 316 66. A summary for the detection of oxygen penetration by microhardness and etching for Ge-doped silicide coatings grown on Ti-22Al-27Nb (22-27) and Ti-20Al-22Nb (20-22) at either 1150°C for 12 hours, 950°C for either 28 hours or 6 hours by three different packs following cyclic oxidation at 500°C, 600°C, 700#C, 800°C, 900°C and 1000°C. (1) Si-Ge/MgF2/SiC, on Ti-22Al-27Nb, 1150°C, 12 hours (2) Si-Ge/MgF2/SiC, on Ti-20Al-22Nb, 950aC, 28 hours (3) Si-Ge/MgF2/SiC, on Ti-20Al-22Nb, 950°C, 6 hours (4) Si-Ge/A1F3/Al203 on Ti-22Al-27Nb, 1150°C, 12 hours (5) Si-Ge/A1F3/Al203 on Ti-20Al-22Nb, 1150aC, 12 hours (6) Si-Ge/A1F3/Al203 on Ti-20Al-22Nb, 950°C, 28 hours (7) Si-Ge/A1F3/Al203 on Ti-20Al-22Nb, 950°C, 6 hours (8) Si-Ge/CuF2/Al203 on Ti-20Al-22Nb, 1150°C, 12 hours (9) Si-Ge/CUF2/Al203 on Ti-20Al-22Nb, 950aC, 28 hours (10) Si-Ge/CuF2/Al203 on Ti-20Al-22Nb, 950°C, 6 hours...... 321 67. The correlation coefficients (R2) and power term (n) for the parabolic and log-log plots shown in Fig. 154 for isothermal oxidation of uncoated Ti-22Al-27Nb ...... 348 68. The correlation coefficients (R2) and power term (n) for the parabolic and log-log plots shown in Fig. 155 for isothermal oxidation of the B-doped silicide coating grown on Ti-22Al-27Nb at 1150°C for 12 hours by a pack composed of 7 wt.% Si, 6% TiB2, 2% MgF2 and A1203 ...... 348 xxxiii 69. The correlation coefficients (R2) and power tern (n) for the parabolic and log-log plots shown in Figs. 156 and 157 for isothermal oxidation of the Ge-doped silicide coatings grown on CP-titanium at 1150°C for 12 hours by a pack comprising either (a) 12 wt.% Si, 6% Ge, 2% MgF2 and Sic or (b) 16% Si, 8% Ge, 2% AlF3 and Al203...... 352 70. The correlation coefficients (R2) and power term (n) for the parabolic and log-log plots shown in Fig. 158 for isothermal oxidation of the Ge-doped silicide coating grown on Ti-22Al-27Nb at 1150°C for 12 hours by a pack composed of 12 wt.% Si, 6% Ge, 2% MgF2 and Sic ...... 354 xxxiv LZBT OF FIGURES FIGURE PAGE 87. Isothermal oxidation at 500-1000eC for 48 hours in air of B-doped silicide coatings grown on CP-titanium by packs comprised of (a) 7 wt.% Si, 6% TiB2, 2% MgF2, Al203 at 1150°C for 12 hours, and (b) 7% Si, 6% TaB2, 2% MgF2, Al203 at 1150°C for 12 hours ...... 212 88. Arrhenius plot of the parabolic rate constant versus inverse temperature compared with oxidation data for Si [156] and TiSi2 [56] for B-doped silicide coatings grown on CP-titanium by packs comprised of (a) 7 wt.% Si, 6% TiB2, 2% MgF2, A1203 at 1150°C for 12 hours, and (b) 7% Si, 6% TaB2, 2% MgF2, A1203 at 1150°C for 12 hours. . . 216 89. SEM micrographs of polished cross-sections of the oxide scales formed on a B-doped silicide coating, which was grown on CP-titanium by a pack comprised of 7 Wt.% Si, 6% TiB2, 2% MgF2, Al2o3 at 1150°C for 12 hours, after isothermal oxidation in air for 48 hours at (a) 900°C and (b) 600°C ...... 222 90. The equilibrium vapor pressures in the B-0 system at 1273 K [143] ...... 225 91. The equilibrium vapor pressures at 1273 K for the (a) Si-0 system and (b) Ti-0 system [143] ...... 228 92. Arrhenius plot of the parabolic rate constant that was corrected for evaporation compared with the measured parabolic rate constant, oxidation data for Si [156] and TiSi, [56]. The B-doped silicide coatings grown on CP-titanium by packs comprising (a) 7 Wt.% Si, 6% TiB2, 2% MgF2, Al203 at 1150°C for 12 hours, and (b) 7% Si, 6% TaB2, 2% MgF2, A1203 at 1150°C for 12 hours...... 230 XXXV 93. Isothermal oxidation of uncoated Ti-22Al-27Nb in air at 600-900°C; (a) weight gain versus time, and (b) Arrhenius plot for the parabolic rate constants compared with reported data for TiSi2[56], pure Ti [10] and other Ti-Al-Nb alloys [161] . . . .232 94. Isothermal oxidation at 500-1000°C for 48 hours in air of a B-doped silicide coating grown on Ti-22Al-27Nb by a pack comprised of 7 wt.% Si, 6% TiB2, 2% MgF2, Al203 at 1150°C for 12 hours. . . 235 95. Arrhenius plot of the parabolic rate constant for a B-doped silicide coating formed on Ti-22Al-27Nb by a pack composed of 7 wt.% si, 6% TiB2, 2% MgF2, AljO, at 1150°C for 12 hours compared with the oxidation data for Si [156] and TiSi2 [56], (a) as-measured and (b) corrected for the evaporation of boron ...... 236 96. The equilibrium vapor pressures at 1273 K for the (a) Nb-0 system and (b) Al-0 system [143] ...... 237 97. SEM micrographs of polished cross-sections of the oxide scales grown on a B-doped silicide coating, which was formed on Ti-22Al-27Nb by a pack comprised of 7 Wt.% Si, 6% TiB2, 2% MgF2, Al20 3 at 1150°C for 12 hours, after isothermal oxidation in air for 48 hours at (a) 900°C and (b) 600°C...... 238 98. Isothermal oxidation at 500-1000°C for 48 hours in air of Ge-doped silicide coatings grown on CP-titanium by packs comprised of (a) 12 wt.% Si, 6% Ge, 2% MgF2, sic at 1150°C for 12 hours, and (b) 16% Si, 8% Ge, 2% A1F3, A1203 at 1150°C for 12 h o u r s ...... 240 99. Arrhenius plot of the parabolic rate constant versus inverse temperature compared with oxidation data for Si [156] and TiSi2 [56] for Ge-doped silicide coatings grown on CP-titanium by packs comprised of (a) 12 wt.% Si, 6% Ge, 2% MgF2, Sic at 1150°c for 12 hours, and (b) 16% Si, 8% Ge, 2% A1F3, A1203 at 1150°C for 12 hou r s ...... 242 100. Vapor equilibria for Ge-0 at 1273K [143]...... 245 xxxvi 101. Arrhenius plot of the parabolic rate constant for Ge-doped silicide coatings with a correction for evaporation compared with the measured rates and oxidation data for Si [156] and TiSi2 [56], The coatings were grown on CP-titanium by packs comprised of (a) 12 wt.% Si, 6% Ge, 2% MgF2, Sic at 1150*C for 12 hours, and (b) 16% si, 8% Ge, 2% A1F3, A1203 at 1150°C for 12 hours ...... 249 102. Isothermal oxidation for 48 hours in air of Ge-doped silicide coating grown on Ti-22Al-27Nb by a pack comprised of 12 wt.% si, 6% Ge, 2% MgF2, Sic at 1150°C for 12 hours, (a) weight change versus time at 500-l000°C and (b) Arrhenius plot of the parabolic rate constants compared with the measured rates and oxidation data for Si [156] and TiSi2 [56] ...... 250 103. SEM micrographs of polished cross-sections of the oxide scales formed on a Ge-doped silicide coating, which was grown on Ti-22Al-27Nb by a pack comprised of 12 wt.% Si, 6% Ge, 2% MgF2, Sic at 1150»C for 12 hours, after isothermal oxidation in air for 48 hours at (a) 900°C and (b)60 0 ° C ...... 251 104. Comparison of Ge-doped and undoped molybdenuro- silicide coatings grown by NaF- and MgF2-activated packs at 500°C, (a) Isothermal oxidation for 2500 hours, and (b) Cyclic oxidation ...... 254 105. SEM micrograph of the surface of a Ge-doped molybdenum-silicide coating after 2500 hours of isothermal oxidation at 500°c, which shows a dual layer oxide scale ...... 255 106. Weight change versus time for the isothermal oxidation at 500°C of a MoSi, heating element and coatings grown by a MnF2-, AIF3- or CuF2-activated p a c k ...... 257 107. Isothermal oxidation at 500°C of undoped and Ge-doped Mo-silicide coatings grown by a NaF-activated pack, with the byproduct salt-oxide layers sanded off some of the coatings, (a) oxidation kinetics and (b) optical photograph after 2500 hours of oxidation [a] sanded prior to oxidation with the olive-green pest product marked as "a", and [b] not sanded, showing no posting product observed ...... 259 xxxvii 108. Weight change versus time for Isothermal oxidation at 500°C of undoped and Ge-doped molybdenum- silicide coatings grown by a NgF2-activated pack that were either untreated or treated with NaF prior to oxidation, (a) unsanded, and (b) sanded to remove byproduct layer before oxidation or NaF treatment ...... 260 109. Optical photograph of molybdenum-silicide coatings grown by a MgF2 activated pack following 500 hours of isothermal oxidation at 500°C in air, (a) Untreated, with massive formation of the olive- green pesting product marked as "a", and (b) treated with NaF prior to oxidation with no pesting product ...... 261 110. Isothermal oxidation at 500®C of a MoSi, heating element that was untreated and treated with NaF prior to oxidation, (a) oxidation kinetics and (b) optical photograph of the untreated (marked "a”) and treated (marked t,b") MoSi2 heating element following 200 hours of isothermal oxidation at 500°c...... 265 111. Isothermal oxidation at 500®C in air for the high- purity MoSi, in the untreated condition and treated with NaF prior to oxidation...... 266 112. Isothermal oxidation at 500°c in air of MoSi2 heating elements in the untreated condition, and (а) treated with NaF and NaCl, and (b) NaBr and Nal prior to oxidation...... 267 113. Cyclic oxidation at 800°C of six different B-doped silicide coatings grown on CP-titanium by packs comprised of : (1) Si-TiB2/MgF2/Al203 [1.7 /«*], (2) Si-TaB2/MgF2/Al203 [6.2 pm], (3) Si-TiB2/AlF3/Al203 [2.6 p m ), (4) Si-TaB2/AlF3/Al203 [3.9 pa] t (5) Si-TiB2/CuF2/Al203 [2.8 Jim], (б) si-TaB2/cuF2/Al2o3 [6.75 /jm], with the thickness for tne TiB2 layers given in brackets . . 272 114. Microhardness profiles that begin at the scale/metal interface on uncoated CP-titanium following 200 cycles at 800°C ...... 274 xxxviii 115. Microhardness profiles that start from the coating/substrate Interface for the B-doped silicide coating on CP-titanium, (a) in the as-coated condition, and (b) following 200 oxidation cycles at 800°C...... 275 116. Optical micrographs of B-doped silicide coatings following 200 oxidation cycles at 800°C, (a) cracks in a coating that was grown by a pack comprised of Si-TiB2/MgF2/Al203, and (b) a crack that penetrated tne substrate for a coating grown by a pack comprised of Si-TaB2/MgF2/Al203. . 277 117. Cyclic oxidation at 875®C for 200 cycles of B-doped silicide coatings grown on CP-titanium at 1150°C for 12 hours by packs comprised of (1) Si-TiB2/MgF2/Al203 or (2) Si-T1B2/CUF2/Al203 ...... 280 118. Cyclic oxidation kinetics for undoped silicide, B-doped silicide and uncoated CP-titanium formed by packs treatments comprised of (1) Si/A1F3/Al203 at 1150*C for 12 hours, (2) Si-TiB2/AlF3/Al203 at 1150«C for 12 hours, and (3) Si-TaB2/AlF3/Al203 at 1150°C for 12 hours; tested in air at (a) 700°C and (b) 8 0 0 ° C ...... 282 119. Cyclic oxidation kinetics for undoped silicide, B-doped silicide and uncoated CP-titanium formed by packs treatments comprised of (1) Si/CuF2/Al203 at 1150°C for 12 hours, (2) Si-TiB2/CuF2/Al203 at 1150°C for 12 hours, (3) Si TiB2/CuF2/Al203 at 950°C for 28 hours, (4) Si-TiB2/CuF2/Al203 at 950°C for 6 hours, and (5) Si-TaB2/CuF2/Al203 at 1150°C for 12 hours; tested in air at (a) 700°C and (b) 8 0 0 ° C ...... 283 120. Cyclic oxidation kinetics for undoped silicide, B-doped silicide and uncoated CP-titanium formed by packs treatments comprised of (1) Si/MgF2/Al203 at 1150°C for 12 hours, (2) Si-TiB2/MgF2/Al203 at 1150°c for 12 hours, (3) Si-TiB2/MgF2/Al203 at 950°C for 28 hours, (4) Si-TiB2/MgF2/Al203 at 950°C for 6 hours, and (5) Si-TaB2/MgF2/Al203 at 1150°C for 12 hours; tested in air at (a) 700°C and (b) 8 0 0 ° C ...... 284 xxxix 121. Cyclic oxidation kinetics at 800°C for B-doped silicide coatings grown on Ti-22Al-27Nb and Ti-20Al-22Nb (a) thick coatings produced at 1150°C for 12 hours and (b) thin coatings produced at 950°C for 6 hours ...... 290 122. Microhardness profile beneath the scale/metal interface of uncoated Ti-22Al-27Nb and Ti-20Al-22Nb alloys following 200 oxidation cycles at 800°C. . . 291 123. Microhardness profiles that start at the coating/alloy interface for the B-doped silicide coating grown on Ti-22Al-27Nb by a pack comprised of Si-TiB2 /MgF2 /Al203 at 1150«C for 12 hours (a) in the as-coated condition and (b) following 200 oxidation cycles at 800°C...... 292 124. Microhardness profiles that begin at the coating/alloy interface for the B-doped silicide coating grown on Ti-20Al-22Nb by a pack comprised of Si-TiB2/MgF2/Al2<>3 at 950°C for 6 hours (a) in the as-coated condition and (b) following 200 oxidation cycles at 800°C ...... 293 125. Cyclic oxidation kinetics for B-doped silicide coatings grown on Ti-22Al-27Nb and Ti-20Al-22Nb by a MgF2-activated pack at either 1150°c for 12 hours, 950°C for 28 hours, or 950°C for 6 hours. Test temperature of (a) 1000°C and (b) 800°C ...... 299 126. Cyclic oxidation kinetics for B-doped silicide coatings grown on Ti-22Al-27Nb and Ti-20Al-22Nb by a CuF2-activated pack at either 1150°C for 12 hours, 950°c for 28 hours, or 950°c for 6 hours. Test temperature of (a) 1000°C and (b) 800°C ...... 300 127. Cyclic oxidation kinetics at 500°C for B-doped silicide coatings grown on Ti-22Al-27Nb and Ti-20Al-22Nb at either 1150°C for 12 hours, 950°C for 28 hours, or 950°c for 6 hours by either (a) a CuF2-activated pack, or (b) a MgF2-activated p a c k ...... 301 128. Cyclic oxidation at 800°C of three different Ge-doped silicide coatings grown on CP-titanium at 1150°C for 12 hours by packs comprised of: (1) Si-Ge/MgF2/SiC, (2) Si-Ge/A1F3/Al203, and (3) Si-Ge/CuF2/Al203 ...... 303 Xl 129. Cyclic oxidation at 875°C of two different Ge-doped silicide coatings grown on CP-titanium at 1150°C for 12 hours by packs comprised of: (1) Si-Ge/A1F3/Al203, and (2) Si-Ge/CuF2/Al203 . . .309 130. Cyclic oxidation at 800°C of Ge-doped silicide coatingB grown by a MgF2-activated pack on CP-titanium at 1150#c for 12 hours with three different Si to Ge ratios in the powder pack; (2:1), (1:1) and (1:2). Also shown is undoped silicide coating and uncoated CP-titanium...... 309 131. Cyclic oxidation at 800°C of Undoped and Ge-doped silicide coatings grown by an AlF3-activated pack on CP-titanium at 1150°C for 12 hours with three different Si to Ge ratios in the powder pack; (2:1), (1:1) and (1:2)...... 310 132. Cyclic oxidation at 800°C of Undoped and Ge-doped silicide coatings grown by a CuF2-activated pack on CP-titanium at 1150°C for 12 hours with three different Si to Ge ratios in the powder pack; (2:1), (1:1) and (1:2)...... 310 133. Cyclic oxidation of the undoped and Ge-doped silicide coatings grown on CP-titanium by a HgF2- activated pack at 1150°C for 12 hours, 950°C for 28 hours, and 950°C for 6 hours tested at (a) 700°C and (b) 800°C for 200 cycles...... 313 134. Cyclic oxidation of the undoped and Ge-doped silicide coatings grown on CP-titanium formed by an AlF3-activated pack at 1150°C for 12 hours, 950°C for 28 hours, and 950°C for 6 hours tested at (a) 700°C and (b) 800°C for 200cycles ...... 314 135. Cyclic oxidation of the undoped and Ge-doped silicide coatings grown on CP-titanium formed by a CuF2-activated pack at 1150°C for 12 hours, 950 °C for 28 hours, and 950°C for 6 hours tested at (a) 700°C and (b) 800°C for 200 cycles...... 315 136. Coefficient of thermal expansion versus temperature for Ti [164], TiSi2 [163] and Ti-24Al-llNb [10] ...... 319 137. The cyclic oxidation kinetics at 800°C for the Ge-doped silicide coating grown on Ti-22Al-27Nb and Ti-20Al-22Nb by a MgF2-, A1F3- and CuF2-activated pack at 1150°C for 12 hours ...... 319 xli 138. Cyclic oxidation of Ge-doped silicide coatings grown on Ti-22Al-27Nb and Ti-20Al-22Nb by a MgF2- activated pack at either 1150°C for 12 hours, 950°C for 28 hours, or 950°c for 6 hours, and tested at (a) 1000°C and (b) 800°C for 200 cycles . 323 139. Cyclic oxidation of Ge-doped silicide coatings grown on Ti-22Al-27Nb and Ti-20Al-22Nb by an A1F3- activated pack at either 1150°C for 12 hours, 950°c for 28 hours, or 950°c for 6 hours, and tested at (a) 1000°C and (b) 800°C for 200 cycles . 324 140. Cyclic oxidation of Ge-doped silicide coatings grown on Ti-22Al-27Nb and Ti-20Al-22Nb by a CuF2- activated pack at either 1150°C for 12 hours, 950°c for 28 hours, or 950°C for 6 hours, and tested at (a) 1000°C and (b) 800°C for 200 cycles . 325 141. Cyclic oxidation at 500°C for Ge-doped silicide coating grown on Ti-22Al-27Nb and Ti-20Al-22Nb by a MgF2-activated pack at 1150°c for 12 hours, 950°c for 28 hours or 6 hours ...... 327 142. Cyclic oxidation at 500°C for Ge-doped silicide coating grown on Ti-22Al-27Nb and Ti-20Al-22Nb by an AlF3-activated pack at 1150°C for 12 hours, 950°C for 28 hours or 6 hours...... 327 143. Cyclic oxidation at 500°C for Ge-doped silicide coating grown on Ti-22Al-27Nb and Ti-20Al-22Nb by an CuF,-activated pack at 1150°c for 12 hours, 950°C for 28 hours or 6 hours ...... 328 144. Optical photograph of B- and Ge-doped silicide coatings grown on Ti-22Al-23Nb fatigue specimens at 950°C for 6 hours...... 328 145. Undoped silicide coatings grown on CP-titanium at 1150°C for 12 hours by packs composed of (a) 10 wt.% Si, 2% A1F3 and A1203, and (b) 10% si, 2% MgF2 and A1203...... 336 146. Undoped silicide coating that was sanded prior to XRD to reveal the inner layers (a) TiSi layer and (b) Ti5Si4 and Ti5Si3 with some TiSi...... 337 147. Boron-doped silicide coatings grown on CP-titanium at 1150°C for 12 hours by packs composed of (a) 7 wt.%Si, 6% TiB,, 2% MgF, and A1203, (b) 7% Si, 6% TaB2, 2% MgF2 ana A1203, and (C) 7% Si, 6% CrB2, 2% MgF2 and A120 3 ...... 338 xlii 148. B-doped silicide coating formed by a pack composed of 7 Wt.% Si, 6% TiB2, 2% MgF2 and Al203 (a) on Ti-22Al-27Nb at 1150°C for 12 nours and (b) on Ti-20Al-22Nb at 950°C for 6 hours...... 340 149. Germanium-doped silicide coatings grown on CP-titanium at 1150°C for 12 hours by packs comprised of (a) 12 wt.% Si, 6% Ge,2% MgF2 and SiC, and (b) 16% Si, 8% Ge, 2% A1F3ana A1203 . 341 150. Ge-doped silicide coatings grown at 1150°C for 12 hours by a pack comprised of (a) 12 wt.% Si, 6% Ge, 2% MgF2 and SiC on Ti-22Al-27Nb, and (b) 16% Si, 8% Ge, 2% A1F3 and A1203 on Ti-20Al-22Nb...... 342 151. Plot of the thickness ratio of Tisi2 over TiSi + Ti5Si4 + Ti5Si3 + Ti3Si for undoped silicide coatings grown at 1150°C, 1050°C and 950°C by packs comprised of (a) 10 wt.% Si, 2% CuF2 and A1203, (b) 10% Si, 2% A1F3 and A1203, and (c) 10% si, 2% A1F3 and Al203 343 152. Linear regression fit for the isothermal oxidation data of a B-doped silicide coating grown on CP-titanium at 1150°C for 12 hours by a pack composed of 7 wt.% Si, 6% TiB2, 2% MgF2, and A1203 (a) parabolic plot and (b) log-log p l o t ...... 345 153. Linear regression fit for the isothermal oxidation data of the B-doped silicide coating grown on CP-titanium at 1150#C for 12 hours by a pack comprised of 7 wt.% Si, 6% TaB2, 2% MgF2, and A1203 (a) parabolic plot and (b) log-log p l o t ...... 346 154. Linear regression fit for isothermal oxidation data of uncoated Ti-22Al-27Nb (a) a parabolic plot and (b) a log-log p l o t ...... 347 155. Linear regression fit for the isothermal oxidation data of a B-doped silicide coating grown on Ti-22Al-27Nb by a pack comprised of 7 wt.% Si, 6% TiB2, 2% MgF2 and A1203 (a) parabolic plot and (b) log-log plot ...... 349 156. Linear regression fit to the isothermal oxidation data for a Ge-doped silicide coating grown on CP-titanium at 1150°C for 12 hours by a pack composed of 12 wt.% Si, 6% Ge, 2% MgF2 and SiC (a) parabolic plot and (b) log-log p l o t ...... 350 xliii 157. Linear regression fit to the isothermal oxidation data for a Ge-doped silicide coating grown on CP-titanium at 1150°C for 12 hours by a pack ‘comprised of 16 wt.% Si, 8% Ge, 2% A1F3 and A1203 (a) parabolic plot and (b) log-log plot ...... 351 158. Linear regression fit to the isothermal oxidation data for a Ge-doped silicide coating grown on Ti-22Al-27Nb at 1150°C for 12 hours by a pack comprised of 12 wt.% Si, 6% Ge, 2% MgF2 and Sic (a) parabolic plot and (b) log-log plot...... 353 159. B-doped silicide coating grown on CP-titanium at 1150°C for 12 hours by a pack comprised of 7 wt.% Si, 6% TiB,, 2% MgF2 and A1203 following isothermal oxidation for 48 hours in air at (a) 900°C and (b) 600°C...... 355 160. B-doped silicide coating formed on CP-titanium at 1150°C for 12 hours by a pack composed of 7 wt.% Si, 6% TaB2, 2% MgF2 and A1203 following isothermal oxidation for 48 hours in air at (a) 800°C and (b) 600°C ...... 356 161. Uncoated Ti-22Al-27Nb following isothermal oxidation for 48 hours in air at (a) 900°C and (b) 600°C...... 357 162. B-doped silicide coating grown on Ti-22Al-27Nb at 1150°C for 12 hours by a pack composed of 7 wt.% Si, 6% TiB2, 2% MgF2 and A1203 following isothermal oxidation for 48 hours in air at (a) 900°C and (b) 60 0 ° C ...... 358 163. Ge-doped silicide coating grown on CP-titanium at 1150°c for 12 hours by a pack comprised of 12 wt.% Si, 6% Ge, 2% MgF2 and Sic following isothermal oxidation for 48 hours in air at (a) 900°C and (b) 600°C...... 359 164. Ge-doped silicide coating formed on CP-titanium at 1150°C for 12 hours by a pack composed of 16 wt.% Si, 8% Ge, 2% A1F3 and A1203 following isothermal oxidation for 48 hours in air at (a) 900°C and (b) 600°c...... 360 165. Ge-doped silicide coating grown on Ti-22Al-27Nb at 1150°C for 12 hours after48 hours ofisothermal oxidation in air at (a) 900°C and (b) 600°c .... 361 xliv 166. Cyclic oxidation of B-doped silicide coatings grown on CP-titanium by six different packs at 1150°C for 12 hours with the thickness of the TiB2 layer given in brackets. Coupons were tested at \a) iooo°c, (b) 900°c, (c) 700°c, (d) 600°c, and (e) 500°C...... 362 167. Cyclic oxidation of undoped and B-doped silicide coating grown on CP-titanium by a MgF2-activated pack at 1150°C for 12 hours or 950®C for 6 hours with the coating thickness given in brackets. Tested at (a) 1000°C, (b) 900°C, (c) 600°C, and (d) 500°C...... 365 168. Cyclic oxidation of undoped and B-doped silicide coating grown on CP-titanium by an AlF3-activated pack at 1150°C for 12 hours with the coating thickness given in brackets. Tested at (a) 1000°C, (b) 900°C, (c) 600°C, and (d) 500°C ...... 367 169. Cyclic oxidation of undoped and B-doped silicide coating grown on CP-titanium by a CuF2-activated pack at 1150°c for 12 hours and 950°C for 6 hours with the coating thickness given in brackets. Tested at (a) 1000°C, (b) 900°C, (c) 600°C, and (d) 5 0 0 ° C ...... 369 170. Cyclic oxidation of B-doped silicide coatings grown on Ti-22Al-27Nb and Ti-20Al-22Nb by MgFj-activated packs at 1150°C for 12 hours, with the boron compound used in the pack given on the plot. Tested at (a) 1000°C, (b) 900°C, (c) 700°C, (d) 600°C, and (e)..500°C...... 371 171. Cyclic oxidation of B-doped silicide coatings grown on Ti-22Al-27Nb and Ti-20Al-22Nb by a CuF2- and MgF2-activated packs at 950°c for 6 hours, with the boron compound used in the pack given on the plot. Tested at (a) 1000°C, (b) 900°C, (c) 700°C, (d) 600#c, and (e) 500° c ...... 374 172. Cyclic oxidation of B-doped silicide coatings grown on Ti-22Al-27Nb and Ti-20Al-22Nb by a CuF2- and MgF2-activated packs at 950°c for 28 hours, with the boron compound used in the pack given on the plot. Tested at (a) 1000°C, (b) 800°C, (c) 5 0 0 ° C ...... 377 173. Cyclic oxidation of Ge-doped silicide coatings grown on CP-titanium by an A1F3-, CuF2- and MgF2-activated pack at 1150°C for 12 hours. Tested at (a) 1000°C, (b) 900#C, (C) 700°C, (d) 600°C and (e) 500°C . . . 379 xlv 174. Cyclic oxidation of Ge-doped silicide coatings grown on CP-titanium by a MgF2-activated pack at 1150°C for 12 hours with a si to Ge ratio of either 2:1, 1:1 or 1:2. Tested at (a) 1000°C, (b) 900*C, (c) 700®C, (d) 600°C and (e) 500°C...... 382 175* Cyclic oxidation of Ge-doped silicide coatings grown on CP-titanium by an AlF3-activated pack at 1150°C for 12 hours with a si to Ge ratio of either 2:1, 1:1 or 1:2. Tested at (a) 1000°C, (b) 900°C, (c) 700°C, (d) 600°C and (e) 500°C...... 385 176. Cyclic oxidation of Ge-doped silicide coatings grown on CP-titanium by a CuF2-activated pack at 1150°c for 12 hours with a Si to Ge ratio of either 2:1, 1:1 or 1:2. Tested at (a) 1000#C, (b) 900®C, (c) 700°C, (d) 600°C and (e) 500°C...... 388 177. Cyclic oxidation of Undoped and Ge-doped silicide coatings grown on CP-titanium by a MgF2-activated pack at either 1150°C for 12 hours or 950*C for 6 hours with the coating thickness in brackets. Tested at (a) 900°C, (b) 600°C, (c) 500°C...... 391 178. Cyclic oxidation of Undoped and Ge-doped silicide coatings grown on CP-titanium by an AlF3-activated pack at either 1150°C for 12 hours or 950°C for 6 hours with the coating thickness in brackets. Tested at (a) 900°C, (b) 600°C, (c) 500°C...... 393 179. Cyclic oxidation of Undoped and Ge-doped silicide coatings grown on CP-titanium by a CuF2-activated pack at either 1150°C for 12 hours or 950#C for 6 hours with the coating thickness in brackets. Tested at (a) 900°C, (b) 600°C, (c) 500°C ...... 395 180. Cyclic oxidation of Ge-doped silicide coatings grown on Ti-22Al-27Nb and Ti-20Al-22Nb by a MgF2-, A1F3-, and CuF2-activated pack at 1150°C for 12 hours. Tested at (a) 900°C, (b) 700°C, and (c) 600°C . . . 397 181. Cyclic oxidation of Ge-doped silicide coatings grown on Ti-22Al-27Nb and Ti-20Al-22Nb by a HgF2-activated pack at 950°C for 6 hours. Tested at (a) 900°C, (b) 700°C, and (c) 600°C...... 399 182. Cyclic oxidation of Ge-doped silicide coatings grown on Ti-22Al-27Nb and Ti-20Al-22Nb by an AlF3-activated pack at 950°C for 6 hours. Tested at (a) 900°C, (b) 700°C, and (c) 600°C...... 401 xlvi 183. Cyclic oxidation of Ge-doped silicide coatings grown on Ti-22Al-27Nb and Ti-20Al-22Nb by a CuFa-activated pack at 950°C for 6 hours. Tested at (a) 900#C, (b) 700°C, and (c) ...... 403 xlvii CHAPTER V OXIDATION RESISTANCE 07 BORON- AND GERMANIUM-DOPED SILICIDE COATINGSt RESULTS AND DISCUSSION 5.1 Isothermal Oxidation The isothermal oxidation kinetics for the coated substrates were monitored in predried air using TGA. The parabolic rate constants were determined by a plot of weight change (mg/cm2) versus the square-root of time [134], The activation energies were determined by an Arrhenius plot of the parabolic rate constants. The mechanisms for high- temperature oxidation were determined by analyzing kinetic data and characterizing the oxidized coupons. 5.1.1 Boron-Dopsd Silicide coatings on CP-Titaniua Figure 87a is plot of weight gain versus time at 500- 1000°C for the B-doped silicide coating grown on CP-titanium from a pack composed of 7 wt.% Si, 6% TiB2, 2% MgF2 and A1203, which shows that several of the oxidation experiments were repeated at the same temperature. The most significant weight occurring during the first couple hours of oxidation was followed by slow, parabolic kinetics. Since the oxidation kinetics for TiB2 are much faster than for Tisi2 [41,56,120,121], the rapid initial oxidation of TiB2 that is 210 211 localized at the coating surface is the source of the faster transient oxidation. To calculate the weight gain produced by complete oxidation of the outer TiB2 layer, a uniform, dense, 1 /xm thick layer of TiB2 was assumed for the B-doped silicide coating grown by a MgF2 activator, which does not agree with the discontinuous TiB2 patches shown in Fig. 42 and the average thickness of 1.21 nm reported in Table 26. The total number of moles for the typical coupon surface area of 2 cm2 is 1.3X10~5. If the 1 jim layer of TiB2 is totally converted into oxide by the following reaction, TiB2 (s) + (5/2) 02 (g) = Ti02 (s) + B203(s ) (5.1) then the transient oxidation weight gain is 0.518 mg/cm2, which is larger than the transient weight gain of 0.1 to 0.3 mg/cm2 observed in Fig. 87a. Since the complete oxidation of the 1 tm TiB2 layer should occur within 1-3 hours at 800- 1000°C [41], and since the TiB2 layer actually consisted of discontinuous patches (Fig. 42), the transient oxidation weight gain involves oxidation of TiB2. Figure 87b shows the weight gain versus time at 500- 1000°C for a B-doped silicide coating grown on CP-titanium by a pack composed of 7 wt.% Si, 6% TaB2, 2% MgF2 and A1203, which has a much thicker layer of TiB2 at the surface of the B-doped silicide coating (Table 26). Since the thickness of the outer TiB2 layer is greater, the transient oxidation weight gain is more significant for these coatings (Note the difference in scale). The oxidation kinetics of pure 212 — *500C *— *600C 0— <>700C *— »800C *— *900C “ olOOOC N E \ cn E o a CT I 20.0 30.0 50.0 (a) Time (hours) ■— *5000 *— *600C o— 0 7 0 OC *— *800C a— n1000C 3.00 ~ 2.50 E 2.00 ■ a> E .c *6 O JC** 1.00 o» I 0.50 0.00 0.0 10.0 20.0 30.0 40.0 50.0 (b) Time (hours) Figure 87. Isothermal oxidation at 500-1000°C for 48 hours in air of B-doped silicide coatings grown on CP-titanium by packs comprised of (a) 7 wt.% si, 6% TiB,, 2% MgF,, A1203 at 1150°C for 12 hours, and (b) 7% si, 6% TaB,, 2% MgF2, A120, at 1150°C for 12 hours. 2 J 213 TiB2 are slower at low temperature [41], which explains the low transient weight gain at 600°C and 500°C. Figures 88a and 88b compare the parabolic rate constants £or the B-doped silicide coatings produced by packs comprised of Si-TiB2/MgF2 /Al203 (thin TiB2 layer) and Si-TaB2/MgF2/Al2o3 (thick TiB2 layer), respectively, with published data for the oxidation of pure Si and TiSi2 [56,156]. The error bars are one standard deviation, and* were calculated by assuming that the data fit a normalized T-distribution [142]. Figures 152a and 153b in Appendix C show that a good parabolic fit was obtained for each plot by neglecting the transient oxidation. The power term (n) for the weight gain (dw) were determined for the following equation, dw - (A)tn (5.2) where A is a constant, by log-log plots in Figs. 152b and 153b in Appendix C. The correlation coefficients for the parabolic and log-log plots with the power terms are summarized in Table 52. The correlation coefficients for the parabolic plots are generally 0.95 of larger, which indicates that the kinetics were parabolic for the majority of oxidation experiments. Since the oxidation kinetics for pure TiSi2 and TiB2 were also parabolic [41,56,120,121], these results are not surprising. However, the log-log plots did not precisely reflect parabolic kinetics because the power terms ranged from 0.2 to 0.8. The majority of power terms ranged from 0.25 to 0.4, which indicates that the kinetics were cubic or quadratic. Kofstad [10] warns against using log- log plots to determine the rate equation. The data are compressed with increasing values of dw and t, which tends to suggest linear relationships between log(dw) and log(t). The rate equation derived from these plots may contradict the true rate equation that was determined by fitting the data, i.e., by a parabolic plot [10]. Furthermore, the correlation coefficients demonstrate that the fit for the parabolic plot was better than for a log-log plot. Therefore, parabolic kinetics were assumed for the isothermal oxidation of these coatings. The error bars shown in Figs. 88a and 88b were small for most oxidation experiments, which indicates that the assumption of parabolic kinetics was reasonable. The oxidation rates for the B-doped silicide coating in Fig. 88a are slower than that for pure TiSi2, but are faster than the oxidation of pure Si. However, the oxidation kinetics for the B-doped silicide coating with the thick layer of TiB2 were much faster than either TiSi2 and pure Si. Thus, small additions of TiB2 improve the oxidation resistance of Tisi2. A change in slope is observed at 700- 750°C for each plot, which indicates a change in the mechanism that controls oxidation. The activation energies for the oxidation of the B-doped silicide coatings are 215 reported in Table 53. The correlation coefficients for the Arrhenius plots are also reported in Table 53. A decent fit was observed for the majority of the data, with some notable exceptions. Despite some of the errors in this data, these activation energies agree with the results discussed in the following sections. At temperatures above 700°C the activation energy for oxidation of the B-doped silicide coating with the thin TiB2 layer is approximately the same as the activation energy for the oxidation of pure Si. This indicates that solid-state diffusion through a Si02 layer controls the oxidation of this B-doped silicide coating on CP-titanium. Since the oxidation of TiSi2 is known to result in the exclusive growth of Si02 at temperatures above 700°C £56/120,121], this result is not surprising. The activation energy for the oxidation of the B-doped silicide coating with the thick TiB2 layer at temperatures above 700°C is very low, and not equal to the value for TiB2, TiSi2, pure Ti or Si. This activation energy is equivalent to the value measured for the B-doped silicide coating with the thin TiB2 layer at 700-500°C, which indicates that the oxidation mechanism was similar. The oxidation of Tisi2 at temperatures below 700°C does not result in the exclusive growth of Si02, but the formation of a mixed Si02-Ti02 scale [56,120,121]. Smeltzer et al. [157] devised a model for the oxidation kinetics of pure Ti at 216 to 1000C BOOC 700C 600C 500C r' E u CM TiSi, Thin Film Oxidation °> 10‘1Z- I 10-’N c o *->0) 1 0 *: ccCO £ 10*1S B -d o p e d 0 Oxidation of Pure Si Silicide Coating on CP-TItanlum 1 10*16- _r f » i ' “ i— r 0.0006 0.0008 0.001 0.0012 0 . 0 0 1 4 (a) 1/T [K'1] t/> 1000C 800C 700C 600C S00C t* E B-doped Silicide o> Coating on CP-Ti 4-*c +4to W c o u CO cc i-14 o Undoped TiSi. o < x» 'P u r e SI 2 01 a. 0 . 0 0 0 8 0.001 0.0012 0 . 0 0 1 4 (b) 1/T [K ] Figure 88. Arrhenius plot of the parabolic rate constant versus inverse temperature compared with oxidation data for si [156] and Tisi2 [56] for B-doped silicide coatings grown on CP-titanium by packs comprised of (a) 7 wt.% Si, 6% TiB2, 2% MgF2, A1203 at 1150°C for 12 hours, and (b) 7% Si, 6% TaB2, 2% MgF2, A1203 at 1150°C for 12 hours. 217 Table 52. The correlation coefficients (R2) determined from the linear regression fit for the a parabolic (dw vs t1/2) and log(dw)-log(t) plot (with the power term (n)) for B-doped silicide coating formed on CP-titanium at 1150°c for 12 hours by packs composed of (a) 7 wt.% Si, 6% TiB2, 2% MgF2, A1203, (b) 7% Si, 6% TaB2, 2% MgF2, Al203. Temp. Parabolic log-log log-1 [•C] (R2) 1000 0.912 0.881 0.271 1000 0.935 0.919 0.279 900 0.992 0.978 0.288 900 0.961 0.962 0.468 800 0.949 0.925 0.679 800 0.893 0.954 0.315 800 0.955 0.901 0.797 700 0.949 0.871 0.414 700 0.954 0.914 0.501 600 0.997 0.979 0.811 600 0.991 0.986 0.587 500 0.891 0.906 0.325 500 0.918 0.821 0.349 Si-TaB2/MgF2 1000 0.982 0.957 0.391 800 0.953 0.874 0.391 700 0.965 0.962 0.352 600 0.986 0.964 0.272 500 0.989 0.994 0.314 218 Table S3. Activation energies (Q[J/mol]) determined for the isothermal oxidation of the B-doped silicide coatings grown on CP-titanium and Ti-22Al-27Nb at 1150°C by packs composed of (a) 7 wt.% si, 6% TiB2, 2% MgF2, A1203, (b) 7% Si, 6% TaB2, 2% MgF2, A12o 3, and reported data for the oxidation of Si, TiSi2, Ti, Nb and TiB2 [10,41,56,156]. The correlation coefficients (R2) for the fit to the Arrhenius plot are also reported, with the proposed mechanism for oxidation. Si-TiB2/ Si-TaB,/ Si-TiB2/ MgF2 (a) MgF2 (b) MgF2 (a) Q [J/mol] 700-1000°C 114,000 18,000 118,000 R 2 700-1000°C 0.895 0.999 0.671 Mechanism Solid-state TiO, Solid-state 700-1000°C Diffusion Graln- Diffusion in Si02 boundary in Sio2 diffusion Q [J/mol] 500-700°C 17,700 172,000 13,600 R2 500-700°C 0.480 0.971 0.697 Mechanism TiO, G.B. TiB, TiO, G.B. 500-700°C Diffusion Oxidation Diffusion Activation Energy for Oxidation of Pure Materials Q [J/mole] Si [156] 119,300 TiSi2 [56] 147,000 TiB2 [41] T > 950*C 125,000 T < 950°C 195,000 Ti [10] 208,000 Nb [10] 117,000 219 600°C to 300®C that was based upon solid-state diffusion along Ti02 grain boundaries, and used an exponential decay term to account for grain growth of Ti02 during an extended period of oxidation. Smeltzer et al. [157] calculated an activation energy of 36,400 J/mole for the low-temperature oxidation of Ti, which was associated with grain-boundary diffusion in the Ti02 scale as the rate-limiting step and grain growth that reduces the number of paths during oxidation. However, the compact Ti02 scale that forms on TiSi2 contains Si02 and B203, which may have doped Ti02 and produced a further change in the mechanism for low- temperature oxidation. Therefore, the low activation energy for the low-temperature oxidation of the B-doped silicide coating with the thin layer of TiB2 probably results from solid-state grain boundary diffusion in Ti02 which contains Si02 and B203. The thick TiB2 layer will be completely converted into Ti02 and B203 during high-temperature oxidation in a couple of hours at 700-1000°C by Eq. (5.1) [41], which is followed by oxidation of the Tisi2 layer. The similarity in activation energy for the low-temperature oxidation of the B-doped silicide coating with the thin TiB2 layer indicates that the oxidation mechanism was the same. Thus, the rate- limiting step for the high-temperature oxidation of this coating was also solid-state grain boundary diffusion through a Tio2 scale that contains sio2 and B203. The activation energy for the oxidation of the B-doped silicide coating with the thick TiB2 layer in the temperature range of 700-500*C is approximately equal to the oxidation activation energy of pure TiB2. From the reported oxidation kinetics for TiB2 [41], the thick TiB2 layer will not be completely consumed during the 48-hour period of oxidation at temperatures below 700°C, which explains why the activation energy for low-temperature oxidation is similar to TiB2. At temperatures above 700°c, the thick TiB2 layer is quickly consumed, and the solid-state diffusional growth of the oxide scale was probably controlled by a grain boundary mechanism through the thick Ti02 layer. Since the low-temperature oxidation of TiSi2 results in the growth of a mixed Ti02-Si02 scale, this mechanism also describes the low-temperature oxidation of the B-doped coating with the thin layer of TiB2. The similarity in activation energy and oxide morphology support this proposed theory. Further investigation is required to prove this proposed mechanism. X-ray diffraction of the B-doped silicide coatings with the thin TiB2 layer following oxidation detected Ti02 (JCPDS# 21-1276) and TiSi2 (Appendix D, Figs. 159a and 159b). Since Sio2 and B2C>3 are amorphous, they were not detected by XRD. Despite the fact that the oxide scale grown by low-temperature oxidation was thinner than the scale formed by high-temperature oxidation, the relative intensity of Ti02 detected by XRD was almost the same. 221 Furthermore, the most intense XRD peak for each pattern was TiSi2, which demonstrates that the whole oxide scale was analyzed by the X-ray beam. This observation indicates that Ti02 comprised a greater fraction of the oxide scale formed by low-temperature oxidation, in agreement with the reported results for low-temperature oxidation of TiSi2 [56,120,121]. The polished cross-sections of the oxide scales for the B-doped silicide coating with the thin TiB2 layer were examined by SEM/EDS following 48 hours of isothermal oxidation at 900°C and 600°C (Figs. 89a and 89b). The high- temperature oxidation of the B-doped silicide coating resulted in a dual-layer oxide scale; (1) outer layer of Tio2 with some Si02 and B203, and (2) thin layer of Si-rich oxide adjacent to the silicide coatings. Although the composition of the inner layer was not pure Si02, these qualitative results indicate that a Si02 layer forms at the TiSi2/oxide interface, in agreement with the thermodynamic stability calculations [18,19] and the oxidation kinetics, since Ti02 and Si02 have limited mutual solubility, the titanium that is liberated by the oxidation of TiSi2 segregates to the surface of the oxide film [158]. Furthermore, Ti02 is more stable than Si02 [143], and would segregate to the high oxygen potential at the scale/gas interface [10]. A single-layered oxide film of Ti02 with Si02 and B203 was detected by EDS at the surface of the B- doped silicide oxidized at 600°C (Fig. 89b), which 222 Inner SI02 Layer Outer TIOa-SIOa layer Figure 89. SEM micrographs of polished cross-sections of the oxide scales formed on a B-doped silicide coating, which was grown on CP-titanium by a pack comprised of 7 wt.% si, 6% TiBj, 2% MgF2, A1203 at 1150°C for 12 hours, after isothermal oxidation in air for 48 hours at (a) 900°C and (b) 600°C. 223 corresponds to the proposed mechanism for the low- temperature oxidation of this coating. Since a continuous second layer of Si02 was not resolved for the oxide scale, the rate-limiting step is probably solid-state diffusion in Ti02 along the grain boundaries that are doped with impurities. The phases detected by XRD following oxidation of the B-doped silicide coating with the thick TiB2 layer were Ti02, TiSi2 and TiB2 for the coupon oxidized at 600°C and Ti02 with minor amounts of TiSi2 for the oxidation at 800°C (Appendix D, Figs. 160a and 160b). Examination of polished cross-sections by SEM/EDS analysis after oxidation at 800°C revealed a thick single-layer scale of Ti02 scale with some Si02 and B203. The single-layered scale following oxidation at 600°c was also Ti02 with B203 on top of the remaining TiB2 layer. This observation confirms that the low- temperature oxidation of the B-doped silicide coating with the thick outer layer of TiB2 is effectively the oxidation of TiB2. The high-temperature oxidation involves solid- state diffusion through a Ti02 scale that contains Si02 and B203, in agreement with the proposed mechanism for the low- temperature oxidation of the B-doped silicide with the thin layer of TiB2. The equilibrium vapor pressures of the B-oxides at 1273 K are given in Fig. 90. The vapor pressure of B02 is high enough at the oxygen partial pressure of air (P02 =0.21) to 224 result In weight changes through evaporation. The vapor pressures of the Si-0 and Ti-0 species at 1273K in Figs. 91a and 91b, respectively, are low in air (PQ2 “0.21), and their evaporation is negligible. Thus, the evaporative weight loss for the B-doped silicide coatings was corrected by the following expression [10], dw/dt = (kpfC/w) - kov (5.3) where the rate constant determined on the left side of Eg. (5.3) is the measured value, kp|C is the parabolic rate constant corrected for the evaporation, and kBV is the evaporation rate constant for B203. Since thermodynamic data for ternary Si02-B203-Ti02 solutions are not known, the evaporation of pure B203 is considered. Since B203 will dissolve into Si02, this calculated rate is an extreme value for the true evaporation kinetics, which is the maximum rate of evaporation. The evaporation of B203 as B02 vapor occurs by the following reaction: B20 3 (c) + (l/2)02 (g) « <2)B02 (g) (5.4) The equilibrium constant (Kg) for Eg. (5.4) is, Kq - (PBo2)2/(Po2>1/2 (5 *5> Since the partial pressure of oxygen is at least five orders of magnitude higher than B02, there is enough oxygen for completion of reaction (5.4). The rate-limiting step for the evaporation of B203 is assumed to be gas-phase diffusion of B02 through a stagnant boundary layer adjacent to the 225 B(s) Stable -S.. PB0. E o CL BO O -20 -23 -33 -30 -20 -10 -3 0 L°g po2 (atm) Figure 90. The equilibrium vapor pressures in the B-0 system at 1273 K [143]. substrate, in accordance with the model of Graham and Davis for the evaporation of Cr203 by Cr03 (Eq. 1.1) [17]. The flux of B02 (J(B02)) can be written in terms of a Fick's first law expression, b R T (5.6) where PB02 is the equilibrium vapor pressure at the substrate, PB02' is the vapor pressure outside the boundary layer, which is assumed to be zero, 6 is the width of the boundary layer, DB02 is the gas-phase diffusion coefficient, T is the absolute temperature and R is the gas constant. The evaporation of l mole of B203 produces 2 moles of B02, which gives the following expression for the evaporation 226 rate (kov): J(B02) = 2J(B20 3) = 2kev/MB203 (5.7) where MB203 is the molecular weight of B203. Substitution of Eqs. (5.5-5.7) gives the expression for the evaporation rate (in units of g/cm2-s): V _ MfyOj DBOj „l/2 nl/4 (5*8) •v 2 6 J?r 0 * Since the oxygen.partial pressure is somuch higher than that for B02, the gas-phase interdiffusion coefficient was determined by assuming that the gas consists of a binary 02- B02 mixture [159], 1.8583*10- (_ L t_L) ( l> "i V (S.9) where a12 is the collisiondiameter, n(1,1) is the omega interval for mass diffusivity, Pt is the total pressure, M the molecular weights and T the absolute temperature. The collision diameter for B02 was estimated from the reported diameter of the molecule [143]. The omega interval for B02 was approximated by using the value for C02, which is similar in size to B02. The data for oxygen are well known, and the values for the omega interval were taken from published tables [159]. Based on the assumption of laminar flow over a flat plate, the width of the stagnant boundary layer is expressed as an integrated averaged over the whole plate [17], 227 (5.10) where L Is the sample length, NSc Is the Schmidt number, N sc “ Vv/°12 (5.11) and NRe is the Reynolds number, NRe = (V)(L)/vv (5.12) where vv is the kinematic viscosity for air, and V is the linear flow rate of the gas. Thus, the evaporation rate of B203 by B02 vapor is calculated by combining Eqs. (5.9) to (5.12) With Eq. (5.8). The evaporation of B203 was negligible at temperatures below 700°C, and the low-temperature oxidation region was not considered. The Arrhenius curves in Figs. 88a and 88b are replotted in Figs. 92a and 92b with the calculated correction for evaporation. The curve corrected for evaporation in Fig 92a falls on the reported kinetics for TiSi2 oxidation, which indicates that B evaporation may explain the appearance of slower oxidation kinetics. However, the evaporation correction shown in Fig. 92a is a maximum rate, which will not occur because of B203 dissolves into Si02 during high-temperature oxidation. Thus, the improvement in the high-temperature oxidation kinetics for Tisi2 by the addition of TiB2 is not an artifact of evaporation. The calculated change in oxidation kinetics for evaporation of the B-doped silicide coating with the thick layer of TiB2 in Fig 92b is insignificant, which shows that evaporation of B203 has little influence on the Figure 91. The equilibrium vapor pressures at 1273 K for for K 1273 at [143). pressures system vapor Ti-0 (b) and equilibrium system The Si-0 (a) the 91. Figure Log P-, (a tm ) 2 Log Pj ( a t m ) 5 2 - -20 -10-. - , r P . - 0 1 - 5 2 - 0 2 5 3 - 0 4 - -* is Stabl Si(s) 5 3 - 0 3 - i Stable b a t S TiO 0 2 - 5 2 - 0 3 - 5 2 - g o L g o L SiO -20 SiO P P q t Oo Stable l b a t S o iO S s u o e itr V q ) m t o ( 2 ) m t a ( 2 5 1 - Oo le b a t S o iO T 10 -1 10 -1 TiO -5 -5 0 0 228 229 oxidation kinetics. One assumption of the evaporation model is the absence of water vapor in the air, which is reasonable for pre-dried air from a gas cylinder, where the partial pressure of H20 vapor is on the order of 10~7 or less [160]. At higher concentrations of water vapor, the evaporation of B203 to create HBO vapor would occur by the following reaction, B203 (s ) + H20(g) - (2)HB0(g) + 02 (5.13) Computer-assisted thermodynamic calculations have shown that the HBO vapor is so stable that water vapor is completely converted to HBO, i.e. for P(H20)»1X10-3 atm., P(HBO)»2X10“3 atm. Oxidation of the B-doped silicide coating with the thin TiB2 layer in air (P(H20)-1X10~3) should result in rapid evaporation of boron from the oxide scale. However, boron was observed in the oxide scale after 200 hours of oxidation in air at 1000°C, which proves that B203 is dissolved into Si02. Therefore, the evaporation of boron does not limit the life of the B-doped silicide coatings. 5.1.2 Uncoated Ti-22Al-27Nb Since oxidation data are not available to make a general comparison, uncoated Ti-22Al-27Nb substrates were oxidized at 600-900°C. Figures 93a and 93b are the plots of weight change versus time and Arrhenius plots for the parabolic rate constants compared with reported results for Ti, TiSi2, respectively. The oxidation kinetics are clearly parabolic for the uncoated alloy (Appendix C, Fig. 154 and 230 c/> 1000C 800C 700C 600C S00C r1 ^ . 1 , l!, ------A ------r ! r E o r>i *12 TiSI2 Thin Film Oxidation O) 10 B -d o p e d c ra Silicide Coating w 10‘1S-J Evaporation c o C o rre c tio n u B-doped >' T \T. *-> 1000C 800C 700C 600C 500C (/> r' E u B-doped Silicide Coating on CP-Ti O) 1 0 ' 1 1 *1 Evaporation C o r r e c tio n c 12 *->ra 10* tn c B-doped Silicide o Coating on CP-Ti CJ 0) 10*’N *-» 03 CC 10*l4iJ Undoped TiSi2 O X) P u r e Si 0) 15 is 10 “ I 1 I I I I | ~T" II I | 11 I I I [--1— Cl. O.OOOB 0.001 0.0012 0.0014 (b) 1/T [K*1] Figure 92. Arrhenius plot of the parabolic rate constant that was corrected for evaporation compared with the measured parabolic rate constant, oxidation data for Si [156] and TiSi2 [56], The B-doped silicide coatings grown on CP-titanium by packs comprising (a) 7 wt.% Si, 6% TiB2, 2% MgF2, A1203 at 1150°C for 12 hours, and (b) 7% Si, 6% TaB2, 2% MgF2, A1203 at 1150°c for 12 hours. 231 Table 67}. The oxidation mechanism for the uncoated alloy * did not change at low temperature and the activation energy is reported in Table 54. The oxide scale detected by XRD was mostly Ti02 with some Nb2o5, and minor a-Al203 (Appendix D, Fig. 161). The oxidation kinetics for uncoated Ti-22Al-27Nb were slower than for pure Ti or Nb, but much faster than for A1203 forming alloys [10]. Furthermore, comparison of the oxidation kinetics for the Ti-Al-Nb alloys in Fig. 93b demonstrates that higher Nb additions improved the oxidation resistance. The activation energy in Table 54 lies between the reported values for the oxidation of pure Ti and Nb, but is much smaller than the activation energies reported for the oxidation of the Ti**Al-Nb alloys that contained less Nb [161). Although no ternary oxide phases were detected by XRD, Ti02 and Nb20s have some mutual solubility [158]. The mutual solubility and limited formation of A1203 may explain why the oxidation kinetics for all the Ti-Al-Nb alloys are slower than for either pure Ti or Nb. Furthermore, previous studies on TiAl have demonstrated that Nb additions stabilized A1203 formation [22,31], and larger Nb additions to Ti3Al-based Ti-Al-Nb alloys may slow oxidation by the formation of more A1203. However, a continuous A1203 scale was not observed, as demonstrated by the fast oxidation kinetics. Thus, the fast rate of oxidation and oxygen dissolution for uncoated Ti-22Al-27Nb clearly demonstrate 232 o ©600C A *700C o 08OOC * — *900C OJ 0.9 -* oE \ o» E c t o™ 0.5 -- o Ol 1 10.0 20.0 30.0 40.0 50.0 (a) Time (hours) to 1000C 900C 800C 700C 600C $ ^ ^ o -n-9 ,> 10 A Ti3AI+5%Nb r i rrlO 10 r Ti3Al+11%Nb Pure Ti | io*n ‘A JS S io’12-i I 10’13n Pure TiSi I i c ’S S3 Ti-22Ai-27Nb § io-ls1 < i | i I i i |— i i i— i | i i — r T - J I CL 0.0007 0.0008 0.0009 0.001' 0.0011 0.0012 (b) 1/T [K ] Figure 93. Isothermal oxidation of uncoated Ti-22Al-27Nb in air at 600-900#C; (a) weight gain versus time, and (b) Arrhenius plot for the parabolic rate constants compared with reported data for TiSi, [56], pure Ti [10] and other Ti-Al-Nb alloys [161], 233 the need for a protective coating. 5.1.3 Boron-Doped Silicide coating on Ti-22Al-27Hb Figure 94 shows the weight change versus time for a B- doped silicide coating grown on a Ti-22Al-27Nb alloy by a pack composed of 7 wt.% si, 6% TiB2, 2% MgF2 and A1203 for 12 hours at 1150°C (Fig. 57a). Host of the weight gain occurs during the first hours of transient oxidation, which is followed by slow, parabolic kinetics, in agreement with the results for the B-doped silicide coating on CP-titanium. The kinetics for each oxidation experiment were generally parabolic (Appendix C, Fig. 155 and Table 68). Two mechanisms for oxidation are resolved in the Arrhenius plot of Fig. 95a, and the activation energies are summarized in Table 53. The parabolic rate constant and activation energy for high-temperature oxidation are similar to the oxidation of pure Si, which demonstrates that solid-state diffusion through a Si02 scale is the rate-limiting step for oxidation at 700-l000°C. The oxidation rates are much slower than for pure Tisi2 and the B-doped silicide coating on CP-titanium. The aluminide phases at the surface probably improved the oxidation kinetics for the B-doped silicide coating grown on the Ti-22Al-27Nb alloy. Figure 95b shows the change in the high-temperature oxidation kinetics for boron evaporation, which were determined by the model developed for the coatings on CP- titanium. The vapor equilibria for Nb205 and Al203 at 234 1000°C in Figs 96a and 96b, respectively, demonstrate that the evaporative losses from these oxides are negligible in air. The oxidation rates are well below the values of pure TiSi2 i but are a little higher than the rates for pure Si. No change in activation energy resulted from the calculated evaporation for the B-doped silicide coating. Thus, boron evaporation does not affect the oxidation kinetics. The change in oxidation mechanism in Fig. 95a occurs at 700°C, in agreement with the results for CP-titanium. The oxide phases detected by XRD were Ti02 and some Nb205 with minor a-Al203 peaks (Appendix D, Fig. 162). Figures 97a and 97b show SEM micrographs of cross-sections for the oxide scales at the surface of the B-doped silicide coating following oxidation at 900°C and 600°C, respectively. Boron was detected in each oxide scale, which demonstrates that complete evaporation did not occur. The oxide scale formed at 900°C was dual-layered, with an outer layer of Tio2-sio2 with B203, A1203 and Nb205. The inner layer adjacent to the TiSi2 coating was almost pure Si02, which agrees with the thermodynamic calculations and kinetic results for high- temperature oxidation. A single-layered oxide scale of Ti02-Si02 with B203, A1203 and Nb2Os was observed following low-temperature oxidation, in agreement with the results for CP-titanium. The activation energy for low-temperature oxidation of the B-doped silicide coating on Ti-22Al-27Nb is also the same as CP-titanium, which indicates that grain 235 Table 54. Activation energies (Q[J/mol]) determined for the isothermal oxidation of uncoated Ti-22Al-27Nb and reported data for the oxidation of TiSi2, Nb, Ti, cto-base Ti-Al-Nb alloys and Al^O* forming alloys tlO,56,1613. Q [J/mol] R2 Uncoated Ti-22Al-27Nb 168,000 0.994 TiSi2 [56] 147,000 -- Ti [10] 208,000 -- Nb [10] 117,200 -- A1 203 [161] 266,000-228,000 -- Ti3Al-5% Nb [161] 488,000 -- Ti3Al-ll% Nb [161] 311,000 -- ■— «500C *— *600C o— 070OC *— *800C •— *900C o— °1000C 0.20 'g 0.15 9-0 01 E ^ 0.10 "5 u | 0.05 0.00 0.0 10.0 20.0 30.0 40.0 50.0 Time (hours) Figure 94. Isothermal oxidation at 500-1000°C for 48 hours in air of a B-doped silicide coating grown on Ti-22Al-27Nb by a pack comprised of 7 wt.% Si, 6% TiB2, 2% MgF2, A1203 at 1150°C for 12 hours. 236 (/> 1000C 800C 700C 600C 500C *■' E ■y v v i i — r o TiSI2 Thin Film Oxidation N O) 10‘'N aVt i o03 -j c o o *14 B -d o p e d ro DC 1 •is Silicide Coating O 10 on Ti-22AI-27Nb o ■ Q Oxidation of Pure Si re ,-16 re 10 Q. i 1 1 i 1-1'— 1— r— 1— r~m*— i— r 0.0006 0.0008 0.001 0.0012 0.0014 (a) 1/T IK-1] (/> 1000C 800C 700C 600C 500C Tl I I I f ^ E u MN TiSi2 Thin Film Oxidation O) 10’,z B -d o p e d c Silicide Coating 4-tre 10°3 tn Evaporation c o C o rre c tio n u *-> Figure 95. Arrhenius plot of the parabolic rate constant for a B-doped silicide coating formed on Ti-22Al-27Nb by a pack composed of 7 wt.% Si, 6% TiB2, 2% MgF2, A1203 at 1150°C for 12 hours compared with the oxidation data for Si [156] and TiSij [56], (a) as-measured and (b) corrected for the evaporation of boron. 237 Nb Stable ^ bC NbO - 10 -- NbO -15-- NbO -25 -35 - 3 0 -25 -20 - 1 5 -1 0 - 5 0 (a) L°9 P 02 (°t m ) Al(s) S able - 5 ” -10 D) - 1 5 ” AIO AIO - 2 0 ” -25 - 4 0 - 3 0 -25 -20 -1 0 -5-35-15 0 (b) Log Pq 2 (atm) Figure 96. The equilibrium vapor pressures at 1273 K for the (a) Nb-0 system and (b) Al-0 system [143], 238 Outer TIP,-SIO,, layer Figure 97. SEM micrographs of polished cross-sections of the oxide scales grown on a B-doped silicide coating, which was formed on Ti-22Al-27Nb by a pack comprised of 7 wt.% Si, 6 % TiB2, 2% MgF2, AI2O3 at 1150°C for 12 hours, after isothermal oxidation in air for 48 hours at (a) 900°c and (b) 600°C. 239 boundary diffusion through a Tio2 layer with B203 and Si02 doping was the rate-limiting step. S.1.4 Germanium-Doped silicide Coatings on CF-Titanlum The weight gain versus tine plot for isothemal oxidation at 500-1000®C in air of Ge-doped silicide coatings grown by a pack composed of either 12 wt.% Si, 6% Ge, 2% MgF2 and Sic, or 16% Si, 8% Si, 2% A1F3, and A1203 is shown in Figs. 98a and 98b. The fast transient oxidation for the Ge-doped silicide coating grown by an A1F3-activated pack is followed by slow parabolic kinetics in Fig. 98b. However, the transient oxidation for the Ge-doped silicide coating grown by a MgF2-activated pack was not very significant. Since the surface of the Ge-doped silicide grown by an A1F3- activated pack contains A1F3 and A1203 deposits, the fast oxidation of the salt deposit, Ge and Ti produced the rapid transient oxidation. Following the transient oxidation, the salt and Ge oxides dissolve into a Si02 film, which results in slow kinetics at steady state. An excellent parabolic fit was observed for the majority of oxidation experiments (Appendix c, Figs. 156 and 157 and Table 69). Figures 99a and 99b compare the parabolic rate constants for the Ge-doped silicide coatings grown by MgF2- and AlF3-activated packs with reported oxidation data for pure Si and TiSi2, respectively [56,156]. The oxidation rates of the Ge-doped silicide coatings are generally slower than pure TiSi2, but a little faster than 240 ■ *500C * »600C ©— o700C *— *800C •— *900C o— *M000C 0.70 0.60 § 0.50” cS J, 0.40 c "5 o 0.30- 0.2 0 -- 0.10-• 0.00 0.0 10.0 20.0 30.0 40.0 50.0 Time (hours) ■— "500C *— *600C o— <>700C *— *800C •— *900C o— °1000C 0.25- c ? 0.20 E*% \ E, 0.15 c *5 o 0.10 JZ O l . a * * - * *• QJ 0.05 *woo0oooo c o ^ 0.00 ( b ) 10.0 20.0 30.0 40.0 50.0 Time (hours) Figure 98. Isothermal oxidation at 500-1000°c for 48 hours in air of Ge-doped silicide coatings grown on CP-titanium by packs comprised of (a) 12 wt.% si, 6% Ge, 2t MgF2, Sic at 1150°C for 12 hours, and (b) 16% Si, 8% Ge, 2% A1F3, A1203 at 1150°C for 12 hours. 241 w 1000C 800C 700C 600C 500C ^ t i r i r i r — ♦— 1 ~ uE TiSI2 Thin Film Oxidation (SIO) 10‘12-=J ro *n*13 G e -d o p e d ♦u> ^ 1 0 c Silicide Coating o O on CP-Titanium 0) *-> ca ae y 10’15 O S 3 Oxidation of Pure Si nj 16 Q. 10 0.0006 0.0008 0.001 0.0012 0.0014 (a) 1/T [K*1] V) r' 1000C 800C 700C 600C 500C uE (SI 12 TiSlj, Thin Film Oxidation *2 10 c CD G e -d o p e d Vt 10‘13n Silicide Coating c o on CP-Titanlum u 4->CD io*’S CD DC U 10',5-J "o Oxidation of Pure Si 2 CD 16 CL 1 0 0.0006 0.0008 0.001 0.0012 0.0014 (b) 1/T [K'1] Figure 99. Arrhenius plot of the parabolic rate constant versus inverse temperature compared with oxidation data for Si [156] and TiSi2 [56] for Ge-doped silicide coatings grown on CP-titanium by packs comprised of (a) 12 wt.% Si, 6% Ge, 2% MgF2, Sic at 1150°C for 12 hours, and (b) 16% Si, 8% Ge, 2% A1F3, A1203 at 1150°C for 12 hours. 242 for the oxidation of pure Si. The rates of isothermal oxidation for the Ge-doped silicide coatings (Fig. 99) are generally much slower than those for the B-doped silicide coatings (Fig. 88). Both plots in Fig. 99 show a reduction in slope at 700-750°C, which indicates a change in the mechanism that controls oxidation. The activation energies for the high-temperature oxidation of the Ge-doped silicide coating grown by an A1F3-activated pack, which are reported in Table 55, are similar to the activation energy for the oxidation of pure Si. Since the rates (Fig. 99b) and activation energy for the oxidation of the Ge-doped silicide coating are approximately equal to the oxidation of pure Si, solid-state diffusion through a Si02 layer is the rate- limiting step. However, the activation energy for the oxidation of the Ge-doped silicide coating grown by a MgF2- activated pack is much closer to the value for the oxidation of pure Ti, which indicates that the oxidation of Ti may have a greater influence on the oxidation kinetics. The activation energies for the low-temperature oxidation of the Ge-doped silicide coatings are low (Table 55), but are approximately equal to the activation energies for the low-temperature oxidation of the B-doped silicide coatings (Table 52). Additionally, the magnitude of the parabolic rate constants for low-temperature oxidation are approximately the same for both the B- and Ge-doped silicide coatings grown on CP-titanium. Since the activation 243 energies and parabolic rate constants for low-temperature oxidation are similar for the B- and Ge-doped silicide coatings, the mechanism of low-temperature oxidation is the same. Therefore, the rate-limiting step for the low- temperature oxidation of the Ge-doped silicide coatings is grain boundary diffusion in a Ti02 scale that is doped with impurities. The oxide phases detected by XRD for the Ge-doped silicide coating grown by the AlF3-activated pack are Tio2 and ct-Al203 (Appendix D, Figs. 163). The A1203 was initially present and forms by the oxidation of the A1F3 layer at the surface of the coating. The presence of A1203 in the oxide scale and loss of fluorine from the oxidation of A1F3 may account for the slower oxidation kinetics observed for the Ge-doped silicide coating formed by the A1F3 activator. Ti02 was the only oxide detected by XRD for the Ge-doped silicide coating formed by a MgF2-activated pack (Appendix D, Fig. 164). A dual-layer oxide scale was observed for both the Ge- doped silicide coatings following oxidation at 900°C in air. The outer layer for the coating formed by the A1F3 activator was found by SEM/EDS to consist of Ti02 with Si02 and Ge02, and isolated regions of A1203, while the thin inner layer was Si02 with larger additions of Ge02 and minor amounts of T i 0 2 . A single layer oxide was observed for both of the Ge- doped silicide coatings after low-temperature oxidation at 244 Table 55. Activation energies (Q[J/mol]) determined for the isothermal oxidation of the Ge-doped silicide coatings grown on CP-titanium and Ti-22Al-27Nb by a pack comprised of either; (a) 12 wt.% Si, 6 % Ge, 2% MgF2, Sic at 1150°C for 12 hours, or (b) 16% Si, 8 % Ge, 2% A1F3, A1203 at 1150°C for 12 hours and reported data for the oxidation of si, Tisi2, Nb, and Ti [10,41,56,156]. The correlation coefficients (R2) for the Arrhenius plot are given with the proposed mechanism of oxidation. AIF3 (b) MgF2 (a) MgF2 (a) CP-Ti CP-Ti 22-27 Q [J/mol] 700-1000°C 88,900 232,000 111,000 R2 700-1000°C 0.941 0.957 0.829 Mechanism Solid-state Solid-state Solid-state 700-1000°C Diffusion Diffusion Diffusion in Si02 in Si02 in Si02 Q [J/mol] 500-700°C 10,800 10,200 -81,900 R2 500-700°C 0.637 0.995 0.993 Mechanism TiO, G.B. TiO, G.B. Coating 500-700°C Diffusion Diffusion Convolutions Activation Energy for Oxidation of Pure Materials Q [J/mole] si [156] 119,300 TiSi2 [56] 147,000 Ti [10] 208,000 Nb [10] 117,000 GeO /•"N E 4 - * - 10 -. O CL c n -15*. 3 Ge(s) Stable Vitreous GeO Stable -25 -35 -30 -25 -1 0 -5 Figure 100. Vapor equilibria for Ge-0 at 1273K [143]. 600°C. The oxide scale for the Ge-doped silicide coating grown by an A1F3 activator consisted of a mixed Ti02-Si02 scale with Al203 and Ge02 additions. The same results were observed for the Ge-doped silicide coating formed by a MgF2- activated pack except that no A1203 was present. Thus, the A1203 additions to the Ge-doped silicide coating that resulted from using an A1F3 activator and A1203 filler produced an improvement in the isothermal oxidation kinetics. Large salt-oxide deposits were not produced by the MgF2-activated pack, and the isothermal oxidation kinetics were faster. The SEM/EDS results indicate that a layer of (Si,Ge}02 was present adjacent to the Ti(Si,Ge)2 coating following high-temperature oxidation, in agreement with thermodynamic stability calculations and the oxidation data. Thus solid- state diffusion through Si02 is the rate-controlling step for high-temperature oxidation. The oxide layer observed following low-temperature oxidation was a single layer of mixed Ti02-Si02, which supports the proposed mechanism of solid-state diffusion in impurity-doped Ti02 along the grain boundaries as the rate-limiting step. The vapor pressures for the oxide phases in the Ge-0 system at 1273 K are shown in Fig. 100. The vapor pressure for GeO is a little higher than B02 at the oxygen pressure of air {P02 =0.21), and evaporation of Ge02 may be important. Since the Ge is dissolved as a solute in the silicide coating and not localized at the surface as a TiB2 layer, modeling the evaporation of Ge02 is more difficult than B 20 3 evaporation. Nonetheless, the evaporation rate of pure Ge02 was calculated as a limiting case for the maximum evaporation rate using Eqs. (5.8- 5.10) for the evaporation of Ge02: Ge02(s) “ GeO(g) + (l/2)02(g) (5.14) The evaporation of 1 mole of Ge02 produces 1 mole of GeO, J(Ge02) - J(GeO) - k^/Mbeoj (5.15) where Mq ^ is the molecular weight of Ge02. Substitution of Eqs. (5.6 and 5.15) gives the expression for the evaporation rate of Ge02 by GeO (in units of g/cm2-s): 247 The maximum possible evaporation rate was found using Eqs. (5.9) and (5.10-5.12). The collision diameter was estimated from the diameter of the GeO molecule [42], and the value for the omega interval was estimated from the critical point for GeO by the approximation of Gaskell [162]. The Arrhenius plots for the Ge-doped silicide coatings grown by MgF2- and AlF3-activated packs with the correction for evaporation are replotted in Figs. 10la and 101b, respectively. The oxidation rates for the Ge-doped silicide coatings are still lower than TiSi2 with the evaporation correction. Thus, evaporation is not much of a factor, and the addition of Ge to the TiSi2 coating has a beneficial effect on the oxidation kinetics for TiSi2. 5 . 1 . 5 Germanium-Doped Silicide coatings on Ti-22Al-27Nb The plot of weight change versus time for a Ge-doped silicide coating grown on Ti-22Al-27Nb by a pack comprised of 12 wt%. Si, 6% Ge, 2% MgF2 and Sic is shown in Fig. 102a. Slow parabolic kinetics are observed following the initial stage of rapid transient oxidation. Since the transient oxidation stage for the Ge-doped silicide grown on CP- titanium by the same pack treatment was insignificant, the aluminide and NbSi2 phases produce the transient oxidation stage observed on the Ti-22Al-27Nb alloy. Parabolic kinetics were generally observed for all the oxidation experiments (Appendix C Fig. 158 and Table 70). Figure 102b shows an Arrhenius plot for the parabolic rate 248 constants with a change in the mechanism of oxidation at 7 0 0 ° C . The rate of oxidation above 7 0 0 ° c is slower than TiSi2 and slightly faster than pure Si. The activation energy for the high-temperature oxidation of the Ge-doped silicide coating (Table 55) is similar to the oxidation of pure Si, which indicates that solid-state diffusion through Si02 is the rate-limiting step. However, the oxidation rate increases at low temperature and a negative activation energy is observed. The oxide scales detected by XRD were Ti02 with some Nb2Os, and the intensities for the Ti02 and Nb2Os peaks were almost as high as for TiSi2 after oxidation at 600 ° C (Appendix D, Fig. 165). The oxide scales observed on the coatings following oxidation are shown in Fig. 103. A dual layer oxide scale was observed following oxidation at 900° C (Fig. 103a): (1) outer layer of Ti02 -sio2 with A1 203, Ge02 and Nb2Os, and (2) inner layer of approximately pure (Si,Ge)02. Thus, solid-state diffusion through Si02 controls the high-temperature oxidation kinetics, in agreement with thermodynamic predictions and the kinetic data. A single-layered oxide scale of Ti02-Si02 with Nb20 5 and Ge02 was observed following the low-temperature oxidation of the Ge-doped silicide coating (Fig. 103b). The reason for the fast oxidation kinetics at low temperature is not clear from the analysis of the oxide scale. However, these Ge-doped silicide coatings were highly convoluted at 249 100X 800C 700C 600C 500C r ' uE TiSf2 Thin Film Oxidation (VI O) c G e - d o p e d (0 ■M Silicide Coating W c Evaporation o C o r r e c t io n (J 0) >14 Ge-doped / ** \ ■ M ra 10 Silicide Coating) cc on CP-Titanium i \ i u io -’S O J 3 ,-16 Oxidation of Pure Si to 10 ■i"-I i" i i i i -T— i— I— r Q. 0.0006 0.0008 0.001 0.0012 0.0014 (a) 1/T[K''] to 1000C 800C 700C 600C S00C r 1 uE (VIO) TISI2 Thtn Film Oxidation i o ° S G e -d o p e d c *->ro Silicide Coating to c & Evaporation o C o r r e c tio n O ft 4-*0) 1 0 ° N G e -d o p e d I ro cc Sllldde Coating u on CP-Tttanlum / o S3 Oxidation of Pure Si e TO 16 CL 10 i i i |1 i" i i [ i i i ] r 1 0.0006 0.0008 0.001 0.0012 0 . 0 0 1 4 (b) 1/T [K'1] Figure 101. Arrhenius plot of the parabolic rate constant for Ge-doped silicide coatings with a correction for evaporation compared with the measured rates and oxidation data for Si [156] and TiSi2 [56]. The coatings were grown on CP-titanium by packs composed of (a) 12 wt.% Si, 6 % Ge, 2% HgF2, Sic at 1150°C for 12 hours, and (b) 16% Si, 8 % Ge, 2% AIF3 , A1 20 3 at 1150°C for 12 hours. 250 ♦— *5000 *— * 600C o— 070OC *— *800C « *900C a— olOOOC 0 . 3 0 j /—X 0 . 2 5 • ■ N E 0 \ 0 . 2 0 ■ 01 E c 0 . 1 5 •• 0 0 £ 0 . 1 0 - ' .? tl S 0 . 0 5 - j 0 . 0 0 1 10.0 20.0 30.0 4 0 . 0 50.0 Time (hours) (/» 1000C 800C 700C 600C 500C v' E * ^ ^ ^ ^""'" V ^ a TiSI2 Thin Film Oxidation s i o - 1^ I 10_1N G e - d o p e d *->• 10'H -J ta Silicide Coating oc on Ti-22AI-27Ntfc ■M 10’,S-J O XI Oxidation of Pure Si ra ,-16 £ 1 0 1— 1 1 j— 1— 1— i— |— i— 1 1 | 1 1 1 '|— 1— 0.0006 0.0008 0.001 0.0012 0.0014 (b) 1/T[K-'] Figure 102. Isothermal oxidation for 48 hours in air of Ge-doped silicide coating grown on Ti-22Al-27Nb by a pack comprised of 12 wt.% si, 6% Ge, 2% MgF2, sic at 1150®C for 12 hours, (a) weight change versus time at 500-l000°c and (b) Arrhenius plot of the parabolic rate constants compared with the measured rates and oxidation data for Si [156] and TiSi2 [56]. 251 Outer TitySlO, layer fa) uter TIOa-S ioa layer (b) Figure 103. SEM micrographs of polished cross-sections of the oxide scales formed on a Ge-doped silicide coating, which was grown on Ti-22Al-27Nb by a pack comprised of 12 wt.% Si, 6% Ge, 2% HgF2, Sic at 1150°C for 12 hours, after isothermal oxidation in air for 48 hours at (a) 900°C and (b) 600°C. 252 the corners and edges, and cracks in the coating were observed. The convolutions observed on the Ge-doped silicide coating were larger than those observed on the B- doped silicide coating (Fig. 59), which did not affect the isothermal oxidation kinetics. The growth kinetics for Si02 are not fast enough to seal or heal these cracks at low temperature, and oxygen probably penetrated into the convolution region. Thus, oxidation of the coating and substrate at the convolutions in the coating produced the fast oxidation kinetics at low temperature. 5.1.6 Low-Tamperature Oxidation Since accelerated oxidation at low temperature is a concern for some silicide compounds (i.e. posting of MoSi2), long-term isothermal oxidation at 500°C was used to test for a rapid rate of oxidation for the B- or Ge-doped TiSi2 coatings. The woight changes measured after 2000 hours of isothermal oxidation at 500°C are recorded in Table 56. The weight changes are small, with demonstrates that accelerated low-temperature oxidation is not a concern for B- and Ge- doped TiSi2 coatings, since Tio2 is not volatile and has a slow growth rate at low temperature [10], a rapid form of low-temperature oxidation is not expected. 5.2 Prevention of Posting for HoSi2 Coatings and Bulk Mo8i2 Posting, or a rapid rate of low-temperature oxidation, is a serious barrier to the wide application of MoSi2. A solution to posting was discovered in this work by studying 253 the low-temperature oxidation of MoSi2 coatings. 5.2.1 Undopad and Ga-Dopad Mo8i3 Diffusion Coatings The isothermal and cyclic oxidation kinetics for MoSi2 and Mo(Si,Ge)2 coatings grown by NaF- or MgF2-activated packs are compared in Figs. 104a and 104b. Posting was observed for both the MoSi2 and Mo(Si,Ge)2 coatings formed by a MgF2-activated pack, as demonstrated by a rapid weight loss resulting from the spalling of a non-adherent, olive- green powder from the coupons. This olive-green powder was collected and identified by X-ray diffraction as Mo03 (JCPDS# 5-508 and 35-609) and Mo902B (JCPDS# 5-441 and 5- 442), which have been reported by other authors as the pasting reaction products [29,30]. No olive-green powder was observed on the surface of Mo(Si,Ge)2 and MoSi2 coatings grown by a NaF-activated pack after 2500 hours of isothermal oxidation or 600 1-hr. oxidation cycles at 500°c, but small amounts of Mo03 and Mo902a were still detected on the surface by XRD. A dual-layer oxide scale is shown in Fig. 105 for a Mo(Si,Ge)2 coating grown by a NaF-activated pack following 2500 hours of isothermal oxidation at 500°C. The composition of the outer layer was determined by EDS to be a Na-rich silicon aluminum oxide, which contains almost no molybdenum, while the inner layer is a molybdenum silicon oxide. The dual-layer scale morphology was also present on undoped MoSi2 coatings grown by a NaF-activated pack. Thus, 254 0.6 MoSL, NaF* Activated McKSi.GeJj, MgFj-Activated v: / MoSij, MgFj-Activated 0.2-1 0.1-1 - 0.1 Mo(Si,Ge)I, NaF* Activate 0 500 1000 1500 2000 2500 Time (hours) 0.150- (—MoS^, MgF2-Activatad 0.125- *5 0 . 0 0 0 Mo(Si.Ga)2, NaF-Activoted -0.025 •• ! Mo(SI,Ge)2 . M gFj-A ctivated -0.050- -Ah 1— ---- 1---- H- + H- 0.0 100.0 200.0 300.0 400.0 500.0 600.0 Figure 104. Comparison of Ge-doped and undoped molybdenum- silicide coatings grown by NaF- and MgF2-activated packs at 500°C, (a) Isothermal oxidation for 2500 hours, and (b) Cyclic oxidation. 255 Table 56. The weight change observed for B-doped and Ge- doped silicide coatings after 2000 hours of isothermal oxidation at 500°c. Weight Change [mg/cm2] CP-titanium Si-TiB2/MgF2/Al203 0.301, 0.0281 Si-TaB2/MgF2/Al203 0.652 Si-Ge/MgF2/SiC 0.105 Si-Ge/A1F3/Al203 -0.704 Si-Ge/Cu F2/A1203 1.331 Ti-22Al-27Nb Si-TiB2/MgF2/Al203 0.190, 0.214 Ti-20Al-22Nb Si-TiB2/MgF2/Al203 0.371 Si-Ge/A1F3/Al203 0.285 Si-Ge/CuF2/Al203 0.451 Outer Na»rlch Oxide Layer tinner Mo-SI Oxide Layer Figure 105. SEM micrograph of the surface of a Ge-doped molybdenum-silicide coating after 2500 hours of isothermal oxidation at 500°c, which shows a dual layer oxide scale. 256 the byproduct Na-rich oxide layer present on the as- coated surface inhibits the occurrence of pesting by forming a sodium-aluminum-silicate scale during low-temperature oxidation. The formation of the sodium-aluminum-silicate passivates MoSi2 by preventing the rapid growth of molybdenum oxides; Fig. 83 shows that the coatings grown by a NaF-activated pack have cracks prior to oxidation, which would have initiated a vigorous pesting reaction if the MoSi2 had been formed otherwise. No sign of pesting was observed in regions that were intentionally precracked with the Vickers microhardness indentor prior to oxidation. The deposited impurity clearly serves to inhibit pesting. The slower oxidation kinetics for the Ge-doped raolybdenum- silicide coating grown by a NaF-activated pack indicate that the addition of Ge has a further beneficial effect on the pesting resistance. The isothermal oxidation kinetics for MoSi2 and Mo(Si,Ge)2 coatings grown by CuF2, A1F3, or MnF2 activators are compared with a MoSi2 heating element in Fig. 106. A comparison of Fig. 104 with Fig. 106 shows that large weight changes are observed for coatings grown by CuF2, A1F3, or MnF2 activators, while a negligible weight change is observed for coatings grown by a NaF activator (note the scales of the plots are different). A considerable amount of non-adherent, olive-green Mo-oxide was observed on the MoSi2 heating element and all other coatings, with the 257 MoSij, AlFj ' ^ ^^MoSij, MnF2 * ¥ | v ' • * ‘ .* , ^M oSL, CuF, ■§, 1- • M • > 3. / / MoSij, CuFj/AljO, & U u ]S) M o (Si,G c)2, AlFj \ I VX ^ Heating Element" Mo(Si.Ge)z, CuF2 MoSij, AIF, t i trp n 11111111111| n 11| i n r p i n p m p n q 500 1000 1500 2000 2500 Time (hours) Figure 106. Weight change versus tine for the isothermal oxidation at 500°c of a HoSi2 heating element and coatings grown by a MnF2-, A1F3- or CuF2-activated pack. 258 exception of coatings grown by a NaF-activated pack. Pesting is recognized in Fig. 106 by either a rapid weight gain or a weight loss upon spelling of the olive green oxide. The composition of the byproduct salt-oxide layer, which results from the use of a specific activator to grow the coating, determines the pesting resistance of the coating. Pesting resistance for MoSi2 was only provided by the byproduct layer that resulted from using a NaF activator. To clarify the advantageous effect of a NaF-activated pack, the byproduct layer was removed mechanically by sanding away 10 /xm of the coating surface. Fig. 107a compares the isothermal oxidation kinetics for sanded and unsanded MoSi2 and Mo(Si,Ge)2 coatings grown by a NaF- activator, whereby larger weight changes are observed for the coatings sanded prior to oxidation. The non-adherent, olive-green oxide that is characteristic of pesting was observed on the surface of coatings sanded before oxidation, absent on unsanded coatings, as seen in Fig. 107b. These results confirm that pesting resistance of the coatings grown by a NaF-activated pack is provided by the byproduct salt-oxide layer. A Na-rich byproduct layer was artificially applied from an aqueous NaF solution to coatings grown by a MgF2- activated pack, which were extremely susceptible to pesting (Fig. 104). The Mg-conraining byproduct deposit on the 259 MoSij, NaF, Sanded' / Mo(Si,Gt)ji NaF, Sandec ■Bb /<>SS '' E, V • ^ \ / - ■ • • /* > ‘ ^ OJJ y « >• § JS MoSi;. NaF • U • t • J3 P • . « * * • Mu £ r-— fSoC sT oe^SS 111111 n v | i 1 1 1 1 11 it 1 1 1 1 1 1 1 1 11 | i 1 1 1 1 1 1 111 m r p 1 11 500 1000 1500 2000 2500 (a) Tim e (hours) □ Figure 107. Isothermal oxidation at 500°C of undoped and Ge- doped Mo-silicide coatings grown by a NaF-activated pack, with the byproduct salt-oxide layers sanded off some of the coatings, (a) oxidation kinetics and (b) optical photograph after 2500 hours of oxidation [a] sanded prior to oxidation with the olive-green pest product marked as "a", and [b] not sanded, showing no pesting product observed. 260 MoSi,, MgF„ Treated ** MoSlj, MgF, 2 f 1 ✓ "6 ■§» Mo(Si.Ge)2, MgF2, Treated e 8) •s, I ■ Mo(S|!gc)2. MgFj ^ MoSi 27 MgF2, T rated 1111111 ii j 11111111111 ■ 11111 ■ 1111 ii 111 ii 11111 jr'iri 600 90 0 1200 1500 (a) Tim e (hours) Mo(Si,Ge),, Sanded, Treated "6 | U) MoSL, Sanded, Treated J=§ U JS * Moo S MoSjj, Sanded Mo(Si,Ge)2, Sanded i n 111111111111111| i > 111 i n 11111111111| 11 r r p T iT 0 300 6 0 0 900 1200 1500 (b) Tim e (hours) Figure 108. Weight change versus time for isothermal oxidation at 500°C of undoped and Ge-doped molybdenum- silicide coatings grown by a MgF2-activated pack that were either untreated or treated with NaF prior to oxidation, (a) unsanded, and (b) sanded to remove byproduct layer before oxidation or NaF treatment. 261 Figure 109. Optical photograph of molybdenum-silicide coatings grown by a MgF2 activated pack following 500 hours of isothermal oxidation at 500°C in air, (a) Untreated, with massive formation of the olive-green pesting product marked as "a", and (b) treated with NaF prior to oxidation with no pesting product. as-coated surface was removed by sanding from some other samples prior application of NaF or oxidation. The Mg- containing byproduct layer was not removed from other coupons, but NaF was applied to determine whether magnesium impurities inherently accelerate pesting. The isothermal oxidation kinetics for MoSi2 and Mo(Si,Ge)2 coatings that were either untreated or treated with a NaF solution are compared for the unsanded and sanded conditions in Figs. 108a and 108b. In both cases, the addition of the NaF salt resulted in slow oxidation kinetics, but a rapid weight 262 change was observed for the untreated coatings. The difference in the oxidation kinetics for the sanded and unsanded samples is not significant. The olive-green pesting product for the untreated coating is clearly shown in Fig. 109, while no sign of the pesting product was observed for coatings treated with NaF prior to oxidation. The addition of a NaF salt layer has provided pesting resistance to MoSi2 and to Mo(Si,Ge)2 coatings grown by a MgF2-activated pack, which are otherwise susceptible to pesting. Thermodynamic calculations for the stability of Si02 on MoSi2 at 500°C using the approach of Rahmel and Spencer [19,143], clearly show that silica is much more stable than the lowest oxide of Mo. Obviously, the competitive growth kinetics for these oxides are the determining factor in pesting. Silica is an extremely slow-growing oxide [10], and the sluggish kinetics for Si02 at low temperature may not enable the exclusive formation of a silica scale. If the faster growing molybdenum oxides form with Sio2, they may result in cracking and disintegration of MoSi2, i.e. pesting. On the other hand, if the growth kinetics for Si02 could be increased by the addition of impurities, then the growth of the molybdenum oxides might be suppressed. The addition of a sodium oxide layer to the surface of MoSi2 may fulfill this requirement by forming a sodium-silicate layer instead of Si02. Sodium ions break up the tightly bound 263 network of Si-0 tetrahedral bonds in Si02, and change the electrical conductivity, viscosity, thermal expansion, diffusion coefficient and mechanical properties [37,38]. This network modifier is commonly used as a flux in glass production [37,38]. The rate of solid-state diffusion in Si02 is increased by the addition of sodium, and the growth kinetics of a sodium-aluminum-silicate scale should be faster than silica [37,38,163]. Thus, the byproduct sodium-oxide layer on MoSi2 and Mo(Si,Ge)2 coatings formed by a NaF-activated pack provides pesting resistance to MoSi2 by the faster growth of a sodium-containing Si02 scale. The relatively rapid growth of the sodium-containing Si02 layer suppresses the formation of excessive molybdenum oxides and passivates MoSi2. Thus, the formation of a Na-containing Si02 is an effective prevention of pesting. Furthermore, the addition of Ge to MoSi2 has been shown to accelerate the growth kinetics of the oxide scale [164], which explains the improvement in the pesting resistance for the Ge-doped MoSi2. However, the most significant improvement was achieved by the presence of a sodium-enriched layer. 5.2.2 Fasting Resistance of Bulk Mosi2 Two types of MoSi2 were used for this study: (l) a used MoSi2 heating element with the oxide layer sanded off; 99% dense MoSi2, produced from powder containing 5-10% sio2 and consolidated by hot pressing [165], and (2) high-purity 264 MoSi2 powder hipped to full (98-99%) density (30]. The previous results showed that pesting resistance was conferred to MoSi2 coatings by NaF, which overwise showed a rapid rate of low-temperature oxidation. The objective of this work was to establish the feasibility of using the NaF salt to prevent the pesting of bulk MoSi2. MoSi2 heating elements were coated with an aqueous solution of either NaBr, NaCl, NaF, Nal or NaN03 in order to clarify the effect of the Na addition. The MoSi2 heating elements were also coated with an organic silicate solution, a boro-silicate, and sodium-silicate solution. The oxidation kinetics at 500°C in air for a MoSi2 heating element are compared to MoSi2 treated with NaF prior to oxidation in Fig. 110a. The pesting of the untreated MoSi2 heating element is shown by the rapid weight loss that results from spalling of an olive-green oxide. This oxide was collected and determined to be Ho03 and Mo9028 by XRD, which is the pesting reaction product £29,30]. A low weight change was observed for the MoSi2 heating element treated with NaF, which demonstrates that the NaF salt can provide pesting resistance to bulk MoSi2. The photograph of untreated and treated MoSi2 heating element after 200 hours of isothermal oxidation at 500°C in Fig. 110b shows the olive-green molybdenum-oxide forming on the untreated sample with no pesting product observed on the treated sample. Thus, the addition of NaF is able to suppress the pesting 265 1.00 0.00 -1.00-- Treated - 2.00" o -4 .0 0 " 5 . 0 0 " - 6.00 " -7.00-- - 8.00 0100 200 300 400 500 Time (hours) Figure 110. Isothermal oxidation at 500°C of a MoSi2 heating element that was untreated and treated with NaF prior to oxidation (a) oxidation kinetics, and (b) optical photograph of the untreated (marked "a") and treated (marked "b") MoSi2 heating element following 200 hours of isothermal oxidation at 500°c. 266 Table 57. The weight change that results from application of the salt layer from an aqueous solution or Si(OC2H5)4, Borax (Na2B407) or sodium silicate solution. Salt w t. Gain t»g/cm2] NaBr 21.07, 18.45, 1.21 NaCl 7.66, 6.03, 4.55, 0.50B NaF 9.38, 8.42, 2.21, 1.91 Nal 4.22, 2.23, 2.22 NaN03 0.764, 0.754, 0.397 Other Solutions Si(OC2H5)4 0.222, 0.097 Borax (Na2B407) 26.55, 8.98 Sodium Silicate 8.45, 7.33, 6.49 Untreated MoSi2 *1s E, & S u § I Treated MoSi III)Jll iii|iiii|iiii|jiii|TnT| 0 200 400 600 800 1000 1200 1400 1600 1800 Time (hours) Figure ill. Isothermal oxidation at 500°c in air for the high-purity MoSi2 in the untreated condition and treated with NaF prior to oxidation. 267 NaF-Treated e *§> E, 1 *; W>4 1 «c NaCl-Treated •SP I 0 200 400 600 800 1000 (a) Time (hours) NaBr-Treated Nal-Treated ■55 £ i: 41 aI ■5) £ 5-1 Untreated I | ITI I | I I I I | TTT |-|-| I I I 200 400 600 800 1000 (b) Time (hours) Figure 112. isothermal oxidation at 500#C in air of MoSi2 heating elements in the untreated condition, and (a) treated with NaF and NaCl, and (b) NaBr and Nal prior to oxidation. 268 reaction on bulk MoSi2. The oxidation kinetics for the high-purity MoSi2 at 500°C are compared with high-purity MoSi2 treated with NaF in Fig. 111. The weight change for the untreated MoSi2 in Fig. Ill demonstrates that the non-adherent, olive-green oxide grew quickly growing (rapid weight gain) and spalled off (weight loss) for these coupons. The low weight changes for the MoSi2 treated with NaF indicates that the olive- green oxide does not form on these coupons. Thus, the NaF provided resistance to the accelerated, low-temperature oxidation of MoSi2 by preventing the rapid growth of Mo oxides, in agreement with the results observed for the MoSi2 diffusion coatings. The NaF addition passivates MoSi2 by forming a fast growing Na-silicate scale during low- temperature oxidation. The weight gain resulting from coating a MoSi2 heating element with a NaF, NaCl, NaBr, NaCl or NaN03 solution, Si(OC2H 5 )4 , Borax or sodium silicate are summarized in Table 57. These weight gains represent the amount of salt coated onto the surface of each coupon. Note that the amount of NaF applied to the MoSi2 used in earlier studies was quite small. The weight gains that result upon isothermal oxidation at 500°c for MoSi2 coated with either NaF or NaCl are compared with untreated MoSi2 in Fig. 112a, while MoSi2 treated with either NaBr and Nal are compared with the untreated MoSi2 in Fig. 112b. The weight change for the untreated MoSi2 is much greater than MoSi2 treated with any « of the Na-salts. Some weight loss is observed for the MoSi2 coupons treated with the Na-salts, but the magnitude for the weight loss does not exceed the weight gain for application of the salt layer (Table 57), which demonstrates that the weight loss is the result of the Na-salt falling off of the MoSi2 coupon and not due to pesting. The olive-green oxide pesting product was not detected following oxidation for any of the coupons treated with a Na-salt, which indicates that any Na-salt will provide pesting resistance to MoSi2. The halide anion is not very important in providing the pesting resistance, although the best pesting resistance was generally produced by the NaF salt. The Na cation is required to form the fast growing Na-silicate that passivates MoSi2 to prevent the pesting reaction, as shown by the results for the MoSi2 diffusion coatings. These results also demonstrate that only a thin layer of Na-salt is required to provide pesting resistance. The weight changes following 500 hours of isothermal oxidation in air for MoSi2 heating elements coated with NaN03, Si(OC2H5)4, Borax (Na2B407) or sodium silicate are summarized in Table 58. A large weight change and massive pesting were observed for NoSi2 coated with NaN03 and Si(OC2H5)4. NaN03 melts below 500°C, and the solution probably dripped off the coupon before a Na-silicate layer could form. Si(OC2H5)4 is an organic precursor for forming 270 Si02, and did not contain the Na additions required to form the fast growing Na-silicate that prevents pesting. The weight changes for the Na-silicate and Borax solutions were quite low, and no pesting product was detected. From the above discussion, Na-silicate is expected to provide pesting resistance, and this observation is no surprise. The results for the Na2B407 solution (Borax) indicate that boron also improves the pesting resistance. Boron forms a fast growing B 20 3 glass at low temperature that may prevent the pesting reaction of MoSi2 in conjunction with the formation of Na-silicate. However, further study is required to resolve which sodium addition (NaF, Nal, sodium silicate or Borax) provides the best pesting resistance. 5.3 Cyolio Oxidation The data are organized as B- and Ge-doped coatings, and subdivided as CP-titanium and Ti-Al-Nb alloy substrates to discuss the cyclic oxidation results. Almost every type of coating was tested at 500°C, 600°C, 700aC, 800°C, 900°C and 1000°C for 200 cycles. The cyclic oxidation kinetics were almost always faster than the isothermal oxidation kinetics. Thermal cycling stresses produce cracks in the silicide coating, which exposes more surface area to oxidation, and produces faster oxidation kinetics. 5.3.1 Boron-Doped silicide coatings on CP-Titanium The six different B-doped silicide coatings used in this study were produced by a coating treatment at 1150°C for 12 hours by packs comprised of: (1) Si-TiB2/MgF2/Al203, (2) Si-TaB2/MgF2/Al203, (3) Si-TiB2/AlF3/Al203, (4) Si- TaB2/AlF3/Al203, (5) Si-TiB2/CuF2/Al203, and (6) Si- TaB2/CuF2/Al203. The average thickness for each layer for the B-doped silicide coatings was reported in Tables 26 to 28. There is some difference in the thicknesses of the silicide layers grown by these packs, but the most significant difference is in the thickness of the outer TiB2 layer. The cyclic oxidation kinetics at 8 0 0 ° c for these six different B-doped silicide coatings and uncoated CP-titanium are shown in Fig. 113, with the thickness of the outer TiB2 layer given in the plot. Large weight gains were observed for the coatings formed with TaB2 in the coating pack, which had the thickest TiB2 layers at the surface. Slow parabolic kinetics were observed for the coatings formed with TiB2 in the coating pack, which had thinner TiB2 layers at the surface of the silicide coating. The silicide coating with the thinnest layer of TiB2 demonstrated the slowest weight gain kinetics at 8 0 0 °C. However, the weight change kinetics for every silicide coating were much slower than uncoated CP-titanium. The inward penetration of oxygen and other contaminants is generally detected by a steep gradient in hardness beneath the scale/titanium interface, and is demonstrated in Fig. 114 by the 250 t m embrittled zone for uncoated CP- titanium after 200 oxidation cycles at 800°C [11,12]. 272 Table 58. The weight gain measured following 500 hours of isothermal oxidation at 500°C for a MoSi2 heating element coated with NaN03, Si{OC2He)4, Borax (Ha2B407) or sodium silicate solution. Wt. Gain [mg/cm2] NaN03 -20.66, -11.75, 20.60 Si(OC2H5)4 13.30, 10.68 Borax (Na2B407) -15.55, -2.61 sodium silicate 1.04, 0.609, -1.68 2.5 W j^Jncoated C P T ^ um^ ^ J d (3.9 um) / (6.75 um) / A * M •a £ 1.5 — —Q I & \ x- ' J, 8 (2.8 um) § ^ 7 ^ (2.6 um) 1 (1.7 um I 1 I 1 I 1 I 1 I 1 I r l 1 I rT 40 80 120 160 Number of Cycles Figure 113. Cyclic oxidation at 800°C of six different B- doped silicide coatings grown on CP-titanium by packs comprised of : (1) Si-TiB2/MgF2/Al203 [1.7 i (2) Si- TaB2/MgF2/Al203 [6.2 pm), (3) Si-T1B2/A1F3/Al203 [2.6 /im], (4) Si-TaB2/AlF3/Al203 [3.9 pm), (5) Si-TlB2/CuF2/Al203 [2.8 pm), and (6) Si-TaB2/CuF2/Al203 [6.75 pm), with the thickness for the TiB2 layers given in brackets. 273 Figure 115 shows hardness profiles from the coating/substrate interface for the B-doped silicide coating formed by a pack composed of 7 wt.% Si, 6% TiB2, 2% MgF2 and A1203, in the as-coated condition (Fig* 115a) and following 200 oxidation cycles at 800°C (Fig. 115b). Since no gradient in the hardness profile nor change in the baseline hardness is observed following 200 oxidation cycles at 800°C, this coating is an effective barrier against the inward penetration of oxygen and other contaminants. Hardness profiles taken in the substrate near deep cracks in the coating (Fig. 116a) did not differ from the profile near an uncracked region. Note the crack in Fig. 116a does not penetrate into the substrate, but is halted by an inner layer of the silicide coating. Then cracking of the brittle silicide coating, which cannot be avoided during cyclic oxidation, does not result in failure of the coating. Additionally, the embrittled zone of uncoated CP-titanium was clearly outlined by using the Kroll etch on the polished cross-section. No embrittled zones were observed for the B- doped silicide coating, in agreement with the microhardness results. The hardness values shown in Figs. 115 and 116 are low for titanium. The use of a hardness standard revealed that the particular machine used in this study gives hardness numbers that are 50-60% lower than the expected value. Nevertheless, the trend in hardness demonstrates that these 274 jp 350.0 E 300.0 o» wu 250.0 0) E 200.0 O z n 150.0 0)n c ■g 100.0 a s ^ O — O w X Q- 50.0 o o * 0.0 0.0 100.0 200.0 300.0 400.0 500.0 600.0 Distance From Scaie/Metal Interface (um) Figure 114. Microhardness profiles that begin at the scale/metal interface on uncoated CP-titaniuro following 200 cycles at 800°C. coatings are protective. The results for the detection of oxygen contamination in CP-titanium substrates coated with all of the B-doped silicide coatings are summarized in Table 59. The only B- doped silicide coatings that were effective barriers against the inward penetration of oxygen for cyclic oxidation at 800°C were grown by packs comprised of either Si- TiB2/MgF2/Al2o3 or Si-TiB2/CuF2/Al203. The general failure mechanism for the B-doped silicide coating with thick layers of TiB2 was rapid oxidation of the TiB2 layer to form large amounts of Ti02i which spalled and caused the underlying silicide coating to crack and fail. Figure 116b shows an 275 CM 300.0 6 E 250.0 + \o> U o 200.0 -O E 3 z n ow c 100.0 "2 o X 50.0 + OQl o c x 0.0 100.0 200.0 300.0 400.0 500.0 600.0 Distance From Coating/M etal Interface (um) 'p 300.0 E "V. 250.0 01 X. fe 200.0 S3 E Z 150.0 c 100.0 a X a 50.0 o o X 0.0 0.0 100.0 200.0 300.0 400.0 Distance From Coating/M etal Interface (um) Figure 115. Microhardness profiles that start from the coating/substrate interface for the B-doped silicide coating on CP-titanium, (a) in the as-coated condition, and (b) following 200 oxidation cycles at 800°C. 276 example of the failure mechanism for a B-doped silicide coating formed by a pack comprised of TaB2/MgF2/Al203, where the oxidized TiB2 layer has spalled off and a crack in the silicide penetrated into the substrate and formed an embrittled zone. The majority of failures in the B-doped silicide coatings occurred at sites that are similar to Fig. 116b. Spalling for the thick layer of oxidized TiB2 is detected in Fig. 113 by a weight loss. Since the B-doped silicide coatings with thick TiB2 layers were susceptible to cracking by spalling of the outer oxide, these coatings were not effective barriers against the inward penetration of oxygen, except at very low temperatures (Table 59). The cyclic oxidation kinetics for the B-doped silicide coatings at 500-1000°C are reported in Fig. 166 of Appendix E, and the results are summarized in the following discussion. Significant weight losses were never observed for any of the silicide coatings at any of the temperatures studied, which indicates that the B-doped silicide coatings did not spall. The weight gain kinetics for the B-doped silicide coating with the thinnest layer of TiB2, which was formed by a pack comprised of si-TiB2/MgF2/Al2C>3 , were generally the slowest at every temperature. Thus, a low amount of boron is desirable for the best coating performance. Large weight gains and non-parabolic kinetics were observed at 900°C and 1000°C, which is above the temperature for the allotropic a to 6 transformation 277 Figure 116. Optical micrographs of B~doped silicide coatings following 200 oxidation cycles at 800°C (a) cracks in coating that was grown by a pack comprised of Si- TiB2/MgF2/Al2 0 3# and (b) a crack that penetrated the substrate for a coating grown by a pack comprised of Si- TaB2/MgF2/Al203. 278 Table 59. A summary for the detection of oxygen penetration by microhardness and etching for B-doped silicide coatings grown on CP-titanium by six different packs at 1150°C for 12 hours following cyclic oxidation at 500*»C, 600°C, 700°C, 800°C, 900®C and 1000*C. The number in brackets is the number of coupons tested: (1) Si-TiB2/MgF2/Al20 3 (2) Si-TaB2/MgF2/Al20 3 (3) Si-TiB2/AlF3/Al20 3 (4) Si-TaB2/AlF3/Al20 3 (5) Si-TiB2/CuF2/Al20 3 (6) Si-TaB2/CuF2/Al20 3 Key: EB - Coating an effective barrier, substrate contamination was not detected. PF - Partial failure of the coating at 1-5 regions with observable substrate contamination. TF « Total failure of coating was substrate contamination. Temp Si-TiB2/MgF2 Si-TaB2/MgF2 Si-TiB2/AlF3 1000°C TF [5] TF [3] TF [2] 900°C TF [4] TF [3] TF [2] 800°C EB [3] TF [3] TF [3] 700°C EB [5] PF [3] PF [2] EB [1] 600°C EB [4] EB [3] EB [3] 500°C EB [4] EB [3] EB [2] Temp Si-TaB2/AlF3 Si-TiB2/CuF2 Si-TaB2/CuF2 1000°C TF [1] TF [3] TF [1] 900°C TF [1] TF [3] TF [1] 800°C TF [2] EB [3] PF [2] 700°C TF [1] EB [3] EB [2] 600°C PF [1] EB [3] EB [2] 500°C EB [1] EB [3] EB [1] 279 temperature (882°C [57]) for titanium. Stresses that result from the allotropic transformation in the substrate created cracks in the silicide coating, and a fast rate of oxidation. The failure of the B-doped silicide coatings at 900°C and 1000°C was characterized by severe embrittlement of the CP-titanium substrate (Table 59). CP-titanium workpieces were coated by the two pack treatments that survived cyclic oxidation at 800°C and tested in cyclic oxidation at 875°C to separate the influence of thermal stresses from the stress produced by the allotropic phase transformation, which occurs at 882°C in the substrate. A rapid rate of weight gain was observed for the B-doped silicide coatings in Fig. 117, and partial failure of each coating and substrate contamination was detected by microhardness and etching. Thus, the high thermal stresses that are produced by cyclic oxidation at higher temperatures are the primary cause of coating failure, independent of the phase change. However, the weight gains that result from cyclic oxidation at 875°C are 2 to 4 times less than those measured at 900°C; this indicates that the stresses from the allotropic phase transformation are also important. For cyclic oxidation at 700-1000°C the weight gains observed for CP-titanium coated with the B-doped silicide are always less than for uncoated samples. However, the weight gain for CP-titanium coated with a B-doped silicide 280 // E ■§» / / / , i—E I s> bo MgF2-Activated Pack (1.7 um)"' 6 § I CuF2-Activated Pack (2.8 um) i-p -rifT m | i i i | i u | i i r p i i*[* 80 120 200 Number of Cycles Figure 117. Cyclic oxidation at 875°c for 200 cycles of B- doped silicide coatings grown on CP-titanium at 1150°C for 12 hours by packs comprised of (1) Si-TiB2/MgF2/Al203 or (2) S i-TiBj/CuF2/A1203. is larger than uncoated substrates at 500°c and 600°c. As previously shown for isothermal oxidation in the preceding section, the low-temperature oxidation rates for the B-doped silicide coatings are faster than the rates extrapolated from high-temperature, because a dual-layered Ti02-Si02 scale is formed instead of Si02. If both the uncoated CP- titanium and the TiSi2 coating form a Ti02 scale, then the weight gain from cyclic oxidation should be comparable. However, the brittle silicide may crack upon thermal cycling, and expose more surface area for oxidation, which 281 produces additional weight gain. Furthermore, the oxidation rate for TiB2 is much faster than for titanium. However, the rate of weight gain decreases to almost zero after about 100 oxidation cycles of low-temperature oxidation for the B- doped silicide coatings, while the uncoated CP-titanium continues to gain weight. Extrapolation of these results to longer times indicates that the weight gain for the uncoated samples exceed those for coated samples. An embrittled zone was observed on the uncoated CP- titanium substrates following low-temperature oxidation. Although the oxide scale on the uncoated CP-titanium substrate could not be resolved after 200 oxidation cycles at 500°, a 5 trn embrittled zone was detected by microhardness. Table 59 shows that no substrate contamination was detected for CP-titanium coated with B- doped silicide coatings. Therefore, the protective coating is required to eliminate contamination of the substrate by oxygen. The cyclic oxidation kinetics at 700°C and 800°C for the B-doped silicide coatings grown by A1F3-, CuF2- or MgF2- activated packs at 1150®C for 12 hours are compared with thinner coatings formed at 950°C and undoped silicide coatings in Figs. 118, 119, and 120. The cyclic oxidation kinetics for the undoped silicide coatings were faster than for the B-doped silicide coatings, which demonstrates the improvement that is provided by the addition of boron. 282 2- I ■*. B-doped (TaB„ 70-60 um) S . i * \ *¥ 1.6 - E, Uncoated CPTi—r ^ * o 60 1.2 J= U «-• 0.8 / .2? B-doped (TiB., 80-70 um] !5 / 0.4 6 - 1 •— undoped Silicide (8 ^Q 0- 11 i I'T'i 11111 | ri 111 111 1111 riT| 1111 11 i | 11 i 0 40 80 120 160 200 (a) Number of Cycles 2.5 - Uncoated CP Ti J *ob ______JS B-dopcd (TaBj, 70-60 um) Mo ■§) B-doped (TiBj, 80-70 um) I Undooed Silicide f80-70 ’I » I 1 I * 80 120 160 (b ) Number of Cycles Figure 1X8. Cyclic oxidation kinetics for undoped silicide, B-doped silicide and uncoated CP-titanium formed by packs treatments comprised of (1) Si/A1F3/Al203 at 1150°C for 12 hours, (2) Si-TiB2/AlF3/Al203 at 1150°C for 12 hours, and (3) si-TaB2/AlF3/Al2o3 at 1150#C for 12 hours; tested in air at (a) 700°C and (b) 800°C. 283 i _____ « ------B-dopwfcTaBj, 70-60 Jim _____ £ £ S Uncoated CP Ti % I .... * X • ■§> X' Undoped Silicide (70*50 p. n) j: « o VO ^ 'B-doped CHBj, 11*8 Jim) • 0.6- B-doped (TIB*, 70;6Q jim)Y^ *§> * r, >• • ~ — * £ 0.3- T T * . T - ^ = : feoped^^, 43-3? Jl 0- l l ~ [ T I T | I I I | t I I | I I 1 | I I I | I I I | I I I p ' l I ) I I I I 80 120 160 200 (a) Number of Cycles Uncoated CP Ti Undoped Silicide (70*50 Jim] "6 •§> B-doped(TiB., 43*35 Jim) £ M B-doped (TaB,, 70*60 Jim)/* V . — *J1 5 6 § W - v*- * ' ' • ^-doped (TiB2, 118 Jim) I B-doped (TiBj, 70-60 jinf) 1 ■ I » 1 1 1 1 I 1 1 1 I 1 I 1 'I-7- 40 80 120 160 200 (b) Number of Cycles Figure 119. Cyclic oxidation kinetics for undoped silicide, B-doped silicide and uncoated CP-titanium formed by packs treatments comprised of (1) Si/CuF2/Al203 at 1150°C for 12 hours, (2) Si-TiB2/CuF2/Al2o3 at 1150°c for 12 hours, (3) Si-TiB2/Cu Fji/A120-j at 950°C for 28 hours, (4) Si- TiB2/CuF2/Al203 at 950°C for 6 hours, and (5) Si- TaB2/CuF2/Al203 at 1150°C for 12 hours; tested in air at (a) 700°C and (b) 800°C. 284 1.2 Uncoated CP Ti H Undoped Silicide (90*80 Jim) 0.8 -^ B-doped (TaB,, 75*65 fim' B i A ■— *- ■— -* — -* 11 u eo N/y ^ ^ “S= I 0.6-1 t x T -- " - '“"""r.-T* / / ^ ♦ • - -B'doped (TiBj, 25-20jim20jxm) 0.4-1 ^ e d (TlBj, 42-35 )im) , : f 0.2-1 B-doped (TiB,, 86*70 |xm i 0 T'i |t~i i | i i i | i i q i i i | i i 11 i i r|~i™i11 fi i 11 111 80 120 (a) Number of Cycles ncoated CP Ti / ! ; » Undoped Silicide (90*80 Jim) / / *§> /■ E -^j^-B-doped (TaBJt 75*65j^ni) ><*!•• 0 g A 'Tar^-tr-*-— —J, U « -> ■a Ii B-doped (tT6 2I^ - 2 0 dlB,. 41 *35 Mm) 1 0.5 -ic __ T ~— * . * B-doped (TiBj, 86-70 Jim) I v" I 1 I 1 i 1 I |1~r i I 1 I 1 1 1 40 80 120 160 200 (b) Number of Cycles Figure 120. Cyclic oxidation kinetics for undoped silicide, B-doped silicide and uncoated CP-titanium formed by packs treatments comprised of (1) Si/MgF2/Al203 at 1150°C for 12 hours, (2) Si-TiB2/MgF2/Al203 at 1150°c for 12 hours, (3) si-TiB2/MgF2/Al2o3 at 950°C for 28 hours, (4) si- TiB2/MgF2/Al2o3 at 950°C for 6 hours, and (5) si- TaB2/MgF2/Al203 at 1150°C for 12 hours; tested in air at (a) 700#C and (b) 800°C. 285 Table 60. A summary for the detection of oxygen penetration by microhardness and etching for the undoped silicide coatings grown on CP-tltanlum by four different packs at 1150°C for 12 hours following cyclic oxidation at 500°C, 600°C, 700°C, 800°C, 900°C and 1000°C. Two coupons were tested at each temperature: (1) Si/MgF2/Al203 (2) Si/A1F,/Al20, (3) Si/CUF2/Al203 (4) Si/MgF2/SiC Key: EB » Coating an effective barrier, substrate contamination was not detected. PF m Partial failure of the coating at 1-5 regions with observable substrate contamination. TF = Total failure of coating was substrate contamination. Temp Si/MgF2/Al203 Si/A1F3/Al203 Si/CUF2/Al203 1000°C TF TFTF 900°C TF TF TF 800°C TFTF TF 700°C TF PF PF 600°C TFPF TF 500#C TFEB PF Temp Si/MgF2/SiC 1000°C TF 900°C TF 800°C TF 700°C PF 600°C PF 500°C EB 286 Substrate contamination was detected for each of the undoped silicide coatings (Table 60), in comparison to no substrate contamination detected for the B-doped silicide coatings formed by CuF2- and MgF2-activated packs. Thus, the boron additions are required for the silicide coatings to prevent the inward penetration of oxygen. The cyclic oxidation kinetics for B-doped silicide coatings with three different thicknesses are also compared in Figs. 119 and 120. A faster rate of weight gain is observed for the thinner coatings. The thinner B-doped silicide coatings were able to prevent the inward penetration of oxygen at 700°C but the thinnest were protective at 600°C and 500°C, as summarized in Table 60. Only the thickest B-doped silicide coatings were protective at 800°C. Thus, thick silicide coatings are required to protect CP-titanium from substrate contamination during long term (200 1-hr. cycles) cyclic oxidation at higher temperatures. The thick B-doped silicide coatings with the best cyclic oxidation resistance at 700-800°C were grown by a MgF2-activated pack at 1150°C for 12 hours. However, the thinner B-doped silicide coatings formed by a CuF2-activated pack were more protective than the thin B-doped silicide coatings produced by a MgF2-activated pack at 700°C. Since the thinner B-doped silicide coatings were processed at 950°C, the more stable MgF2-activator produces a low 287 Table ex. A summary for the detection of oxygen penetration by microhardness and etching for the thin B-doped silicide coatings grown on CP-titanium by two different packs at either 950°C for 6 hours or 950°C for 28 hours following cyclic oxidation at 500°C, 600°C, 700®C, 800°C, 900®C and 1000°C, with the number of coupons tested at each temperature in brackets: (1) Si-TiB2/MgF2/Al203 at 950°C for 6 hours (2) Si-TiB2/CUF2/Al203 at 950#C for 6 hours (3) Si-TiB2/MgF2/Al203 at 950°C for 28 hours (4) Si-TiB2/CuF2/Al203 at 950*C for 28 hours Key: EB « Coating an effective barrier, substrate contamination was not detected. PF « Partial failure of the coating at 1-5 regions with observable substrate contamination. TF - Total failure of coating was substrate contamination. Temp Si-TiB2/MgF2 950°C,6h Si-TiB2/CuF2 950°C,6h 900°C TF [2] TF [2] 800°C TF [2] TF [2] 700°C PF [2) TF [2] 600°C EB [2] EB [2] 500#C EB [2] EB [2] Temp Si-TiB2/MgF2 950°C,28h Si-TiB2/CuF2 950°C,28h 09 TF [2) TF [2] o o e O O O o oo oo PF [1] EB [3] EB [2] 288 gas-phase flux of boron-fluorides and forms less TiB2 at the surface of the silicide coating. The Si02 scale formed during high-temperature oxidation may not contain enough boron to provide an improvement for the coating grown by a MgF2-activated pack at low temperature. Thus, a less stable activator is required to increase the gas-phase flux and deposit sufficient boron at the surface of coatings grown at lower temperatures. The cyclic oxidation kinetics at 500°C, 600°C, 900'C and 1000°C for the B-doped silicide coatings formed by MgF2, A1F3 and CuF2 activators are compared with the thin coatings and undoped silicide coatings in Figs. 167, 168, and 169, of Appendix E. The weight gain kinetics for the B-doped silicide coatings were generally slower than the undoped silicide coatings at all temperatures. Additionally, 500°C was the only temperature at which the undoped silicide coating could prevent substrate contamination (Table 60). Thus, the boron additions provide the cyclic oxidation resistance to the silicide coatingB. 5.3.2 Boron-Doped Silicide coatings on Ti-Al-Nb Alloys The thicknesses for the B-doped silicide coatings tested by cyclic Oxidation were reported in Table 33. The differences between the B-doped silicide coatings grown at 1150°C and 950°c are the thicknesses of coating layers and the width and composition of the interdiffusion zone adjacent to the coating/alloy interface. 289 The cyclic oxidation kinetics at 800°c for the thick and thin B-doped silicide coatings, formed at 1150°C and 950°C, respectively, are shown in Figs. 121a and 121b. The B-doped silicide coatings provide an enormous improvement compared to the uncoated Ti-22Al-27Nb and Ti-20Al-22Nb alloys. The cyclic oxidation rates increase slowly, and the kinetics exceed the rate of isothermal oxidation. The cyclic oxidation kinetics at 500°C, 600°C, 700°C, 900°C, and 1000°C for the thick and thin B-doped silicide coatings on the Ti-Al-Nb alloys are given in Figs. 170, 171 and 172 in Appendix G. Figure 122 shows that a 75-50 n m embrittled zone is observed beneath the scale/metal interface of uncoated Ti- 22Al-27Nb and Ti-20Al-22Nb following 200 oxidation cycles at 800°C. The hardness gradient is very steep in the contaminated zone adjacent to the substrate, which is an excellent site for the nucleation and growth of cracks [12- 14]. Figure 123 shows hardness profiles for the thick B- doped silicide coating on Ti-22Al-27Nb following 200 cycles at 800°C compared to the as-coated condition. The hardness gradient observed in the as-coated condition results from the enrichment of A1 at the coating/alloy interface. A comparison of Fig. 123a with the EDS profile in Fig. 57b shows a direct correlation between Al-enrichment in the interdiffusion zone and increased hardness. Figure 124 shows hardness profiles for the thinnest B-doped silicide 290 1.2 7~~ Uncoated Ti-20Al-22Nb 0.8 Uncoated Ti-22AI-27Nb 0.6 0 .4 - 3 0.2 TiBj, (22-27) 8 0 1 2 0 1 6 0 200 N u m b e r o f C y c le s 0.8 h Uncoated Ti-20Al-22Nb Ts Uncoated Ti-22AJ-27Nb v * . . , . • 4 ^ 7 C u F 2- TiB2 (20-22) ■§) 0.6-3 JS ■ ' " * OX) lit § J3 0 . 4 | U U a*S) j, /•'* MgFj, TiB2 (20-22) MgFj, TiBj (22-27) £ 0.2 CuF?,TaBj (20-22) 0 1— 1 1 I 1 I ' I 1 I 1 I ' 1 1 I 1 I 1 80 120 160 200 (b) Number of Cycles Figure 121. cyclic oxidation kinetics at 800°C for B-doped silicide coatings grown on Ti-22Al-27Nb and Ti-20Al-22Nb (a) thick coatings produced at 1150°c for 12 hours and (b) thin coatings produced at 950°C for 6 hours. Figure 122. Microhardness profile beneath the scale/metal scale/metal 800°C. the at cycles beneath profile oxidation 200 Microhardness following 122. Figure interface of uncoated Ti-22Al-27Nb and Ti-20Al-22Nb alloys alloys Ti-20Al-22Nb and Ti-22Al-27Nb uncoated of interface Knoop Hardness Number [kg/mm ] 700-ar 2 4 6 8 10 2 10 160 140 120 100 80 60 40 20 0 Uncoated Ti-20Al-22Nb Uncoated Uncoated Ti-22Al-27Nb Uncoated Distance [pm] Distance Igj- q _ n — 1 ■ o -= e P 291 z 150.0 m s c 100.0 - "2 o X a 5 0 .0 - o o c («) X 0.0 4- + + 0.0 100.0 200.0 300.0 400.0 500.0 600.0 Distance From Coating/M etal Interface (um) f t * 300.0 E •s. 250.0 + o> l _ 0 200.0 ■ O 3E z 1 5 0 .0 - 100.0 - ■8 o X a 5 0 .0 - o o 100.0 200,0 300.0 400.0 Distance From Coating/M etal Interface (um) Figure 123. Microhardness profiles that start at the coating/alloy interface for the B-doped silicide coating grown on Ti-22Al-27Nb by a pack comprised of si- TiB2/MgF2/Al203 at 1150°C for 12 hours (a) in the as-coated condition and (b) following 200 oxidation cycles at 800°C. 293 250 N el> 2004 1 150-i 1004 a, 504 ■|ii)HiiiHiii|imiiiiiiniii||iii|inUi'rrl 250 g 2004 1504 1 1004 X 504 8 3 “ 11 l l | I I I 1111 ii 11 i i 11 ii 11[ i i 11 | i i n 11 ii 11 m i [ i i ii | i U i | m i | u 1111 i i 1111 ii | ii 1111111 0 20 40 60 80 100 120 140 160 Distance Qun] Figure 124. Microhardness profiles that begin at the coating/alloy interface for the B-doped silicide coating grown on Ti-20Al-22Nb by a pack comprised of Si- TiB2/MgF2/Al203 at 950°C for 6 hours (a) in the as-coated condition and (b) following 200 oxidation cycles at 800°C. 294 coating on Ti-20Al-22Nb in the as-coated condition and after 200 oxidation cycles at 800°C. The hardness gradient that results from the coating treatment at 950°C is much thinner, and the width of this zone is in agreement with the depth of the Al-rich interdiffusion zone shown by the EDS profile in Fig. 60b. No change was observed in either hardness profile following 200 cycles at 800°C, which indicates that the coating protects the alloy from the inward diffusion of oxygen. A summary for the detection of substrate contamination for the B-doped silicide coatings on the Ti- Al-Nb alloys is reported in Table 62. The thick B-doped silicide coatings grown on the Ti- 22Al-27Nb alloy were more protective than the same coating formed on Ti-20Al-22Nb. The thick B-doped silicide coating formed on the Ti-20Al-22Nb alloy did not prevent the inward penetration of oxygen. The B-doped silicide coatings formed on Ti-20Al-22Nb were much thicker (Table 33), and were highly convoluted at the corners, which were the sites of coating failures. However, failure of the thick B-doped silicide coatings on Ti-20Al-22Nb was not observed at the flat regions, which were not convoluted. The B-doped silicide coating grown on Ti-20Al-22Nb by a pack containing TaB2, which produced a thicker layer of TiB2 at the substrate surface, showed a larger weight gain at 800°c, but prevented substrate contamination, since this coating was not as thick (Table 33), the convolutions were not as large. 295 Table 62. A summary for the detection of oxygen penetration by microhardness and etching for the B-doped silicide coatings grown on Ti-22Al-27Nb (22-27) and Ti-20Al-22Nb (20-22) by the pack treatments below following cyclic oxidation at 500-1000°C, with the number of coupons tested at each temperature in brackets: (1) Si-TiB2/MgF2/Al203 on Ti-22Al-27Nb at 1150°C, 12h (2) Si-TiB2/MgF2/Al20, on Ti-20Al-22Nb at 1150°C, 12h (3) Si-TiB2/MgF2/Al203 on Ti-20Al-22Nb at 950°C, 28h (4) Si-TiB2/MgF2/Al203 on Ti-22Al-27Hb at 950°C, 6h (5) Si-TiB2/MgF2/Al203 on Ti-20Al-22Nb at 950°C, 6h (6) Si-TaB2/MgF2/Al203 on Ti-20Al-22Nb at 1150°C, 12h (7) Si-T1B2/CUF2/Al203 on Ti-20Al-22Nb at 950°C, 28h (8) Si-TiB2/CuF2/Al203 on Tl-20Al-22Nb at 950°C, 6h (9) Si-TaB2/CUF2/Al203 on Ti-20Al-22Nb at 950°C, 28h (10) Si-TaB2/CuF2/Al203 on Ti-20Al-22Nb at 950°C, 6h Key: EB - Coating an effective barrier, substrate contamination was not detected. PF - Partial failure of the coating at 1-5 regions with observable substrate contamination. TF - Total failure of coating was substrate contamination. Temp Si-•TiB2/MgF2 Si-TlB2/MgF2 si■TaB2/MgF2 (1) 22 -27 (2) 20'-22 te) 20-22 1000°C EB [4] TF [1] PF [1] 900°C EB [4] PF [1] EB [1] 800°C EB [3] PF [1] EB [1] 700°C EB [4] PF CD PF [1] 600°C EB [2] PF [3] PF [1] PF [1] 500°C PF [4] PF [1) PF tl) Si- TiB.?/MgF2 si■T1B2/CUF2 Si-■TaB2/CuF2 (5) (4) ts) 20'-22 (10) 20-22 1000°C EB [2] PF [1) TF [1] TF [1) 900°C EB [3] EB Cl] TF [1] EB [1] TF [1] 800°C EB [2] EB [1] TF [2) EB [1] 700°C EB [3] EB [1] EB ID 600°C EB [3] EB [2] EB tl] 500°C EB [2] EB [1] EB [2) EB [2] siTiB2/MgF2 si-TiB2/CuF2 Si-■TaB2/CuF2 ts) 20 -22 (7) 20a-22 OJ 20-22 1000°C EB 13 3 EB [2] PF [1] PF [2] 800°C EB [2] EB [2] EB [2] 500°C EB [1] PF [1) EB [1) 296 Additionally, the increased B-doping of Si02 that results from the thicker TiB2 layer provided more effective sealing of cracks in the coating. The trends in cyclic oxidation kinetics observed for the thick B-doped silicide coatings at 800°C were generally also observed at 700°C-1000°C. The oxidation rates for the alloys coated with the B-doped silicide were slower than for the uncoated alloy. The B-doped silicide coating on Ti- 22Al-27Nb was always protective and provided the slowest oxidation rate, but the weight gains for the B-doped silicide with the thickest layer of TiB2 were always the largest. However, large weight losses were observed for cyclic oxidation at 500°C and 600°C due to spalling of the coating at the convoluted corners (Fig. 59). Substrate contamination and an oxide scale was detected by EDS at the corners where failure had occurred. As previously mentioned, the oxide scale that grows on TiSi2 at low- temperature is a mixed Ti02-Si02 scale. The convoluted corners (Fig. 59) probably cracked during cyclic oxidation, and the volume change resulting from Ti02 growth produced further crack growth. Additionally, the NbSi2 precipitates in the silicide coating were also oxidized to further accelerate crack growth at the corners. The coatings were protective at flat regions, which were not convoluted. The cyclic oxidation kinetics for the thinnest B-doped silicide coatings formed by a MgF2-activated pack were the 297 same on Ti-22Al-27Nb and Ti-20Al-22Nb (Fig. 121b). Substrate contamination was not detected for most of the thin B-doped silicide coatings (Table 62), with the exception of the coating failure and oxygen embrittlement observed for the coating grown by a pack comprised of Si- TiB2/CuF2/Al203, which had the thinnest TiB2 layer. The failure of this B-doped silicide coating is indicated in Fig. 121b by the linear weight gain dependence and fast rate for the cyclic oxidation kinetics. The cyclic oxidation kinetics for the thin B-doped silicide coatings at 500°C, 600°C, 700°C, 900°C and 1000°C are shown in Figs. 171 and 172 in Appendix E, and are summarized in the following discussion. The oxidation kinetics for the B-doped silicide coatings were much slower than uncoated alloys at 700-1000°C. However, the cyclic oxidation kinetics for the coated and uncoated substrates were similar at 500°C and 600°C. The low-temperature oxidation mechanism for TiSi2, which produces a mixed Ti02- Si02 scale, and cracking of the silicide coating produced this result, since the thin silicide coatings were not convoluted at the corners, spalling of these coating was not observed. The thinnest B-doped silicide coatings grown by a CuF2- activated pack prevented the inward penetration of oxygen for cyclic oxidation at temperatures up to 700°C and 900°C for packs comprised of either Si-TiB2/CuF2/Al203 or Si- 298 TaB2/CuF2/Al203, respectively. However, B-doped silicide coatings formed by a MgF2-activated pack were generally protective at 500-1000°C, with one partial failure observed at 1000°C. Additionally, slow cyclic oxidation kinetics and protection from substrate contamination was observed at 1000°C for the B-doped silicide coating on a o2-base Ti- 24Al-llNb alloy (Fig. 125a). The biggest difference between the coatings formed by CuF2 and MgF2 activators is the thickness of the coating (Table 33). Thus, a coating treatment of 950°C for 28 hours was used to double the thickness of each coating in an effort to improve the cyclic oxidation resistance at high temperature. Figures 125 and 126 compare the cyclic oxidation kinetics at 1000°C and 800°C for all B-doped silicide coatings formed by MgF2- and CuF2-activated packs, respectively. The oxidation kinetics for the B-doped silicide coatings formed at 950°C for 28 hours are slow, and all of these coatings prevented substrate contamination at 1000°C and 800°C. However, previous results demonstrated that the thicker coatings were susceptible to spalling during low-temperature oxidation. No convolutions were observed at the corners of these coatings, and the cyclic oxidation results in Figs. 127a and 127b demonstrate that no weight loss nor spalling was generally observed for cyclic oxidation at 500°C. Some spalling was observed for the B- doped silicide coating formed by a pack comprised of 299 2.4 UncoatedTi-22Al-27Nb^20 22; TaB^ x 10. 85 ^ m j 2 \ "e (20*22: TiBj, 130-115 Jim) *§> B 1.6*1 O (20-22: TiB2, 51-44 jim )^ , ...... 60 \- - m § j / X- - ' * ' \ (20-22: TiB , 25-18 Jinj) e ( (24-1 l:TiBj, 20-10 Jim) \ A1 • • 0.4*= Jn coated Tt-20A1-22Nb <22-27: TiB,, 130-115 ] 40 80 120 160 (a) Number of Cycles 1.2 L ^Uncoated Ti-20Al-22Nb / / Upcoatedcoated Tl-22Al-27Nb l-l "E “ (20-22: TaB , 110-85 Jim) A x ■§» £ I * (20-22: TiBj, 130-115 Jim) E 0.8-! 4) \ I ■' / (20-22: TiB,, 25-18 Hm) 60 0-6i '/ S (22-27: TiB , 125-110 ji . f .4 -I ,'*'(22-27: T iB „ 2 5 -1 5 s (b) 40 80 120 160 Number of Cycles Figure 125. Cyclic oxidation kinetics for B-doped silicide coatings grown on Ti-22Al-27Nb and Ti-20Al-22Nb by a MgF2- activated pack at either 1150°C for 12 hours, 950°C for 28 hours, or 950°C for 6 hours. Test temperature of (a) 1000°C and (b) 800°C. 40 80 120 160 200 (a) Number of Cycles U ,L- Uncoated Ti-20Al-22Nb ^ • *' .** v- Uncoated Ti-22Al-27Nb-^ (20-22: TaB , 26-2 / X ■§) . -22: TiB,. 15-10 jxtfi) & I1 • ' * •' & h s § 6 f « • " (20-22: TiB,, 47-40 um) S s — — ■ — —■ — ■— “* 20-22: TaB , 15-10 pm) 80 120 160 200 (b) Number of Cycles Figure 126. Cyclic oxidation kinetics for B-doped silicide coatings grown on Ti-22Al-27Nb and Ti-20Al-22Nb by a CuF2- activated pack at either 1150°C for 12 hours, 950°C for 28 hours, or 950°c for 6 hours. Test temperature of (a) 1000°C and (b) 800°C. 301 k • • * - f ' ------(20*22: TaBa , 26-20 Jim) j Uncoaled Ti-22Al-27Nb UncoatedTi-20Al-22Nb 1b £ 3 V u t>o (20-22: T aB ,, 15-10 um) (20-22: TiBJt 15-101 (20-22: TiB., 47-40 pm ). i i~i | i i iy i i i |-i i i | i i i | i i i | i i ,*j . i i | i i i | i i i 0 40 80 120 160 200 (a) Number of Cycles (20-22: TiBj, 51-44 oated Ti-20A1 2 2 N b Uncoated Ti-22Al-27Nb *¥ ■§> £ o ' (22-27: TiB , 125-110 pm) 00 s f s (22-27: TaBj, 110-85 ’I) £ in(20-22: 11 i !■ TiB..i | at 130-115i-1 i'i i ['iuml ri | 11 n ' l T i j r n | i r» | i i 80 120 160 (b) Number of Cycles Figure 127. Cyclic oxidation kinetics at 500°C for B-doped silicide coatings grown on Ti-22Al-27Nb and Ti-20Al-22Nb at either 1150°C for 12 hours, 950°c for 28 hours, or 950°c for 6 hours by either (a) a CuF2-activated pack, or (b) a MgF2- activated pack. 302 Si-TiB2/CuF2/Al203, and substrate contamination was detected. 5.3.3 Germanium-Doped Silioida Coatings on CP-titanium As previously discussed, the amount of Ge addition in the Ge-doped silicide coating is decided by the halide activator and inert filler used in the pack powder. Tables 34, 35 and 36 show that the thicknesses for the three different coatings are not the same. The cyclic oxidation kinetics at 800°C for the Ge-doped silicide coatings grown at 1150#C for 12 hours by a MgF2-, A1F3-, and CuF2-activated pack are compared to uncoated CP-titanium in Fig. 146. The weight gain kinetics for the Ge-doped silicide coatings are a vast improvement over uncoated CP-titanium. The cyclic oxidation weight gains were slowest for the Ge-doped silicide coating formed by a MgF2 activator, which contained the highest Ge addition. The cyclic oxidation kinetics for the coating formed by the MgF2-activated pack are linear, but the kinetics for the coatings formed by the A1F3- and cuF2-activated packs are approximately parabolic. Substrate contamination was observed for the Ge-doped silicide coatings formed by the MgF2 activator, and Table 63 shows that the coatings formed by the A1F3 and CuF2 activators were protective. Thus, the cyclic oxidation kinetics predict whether the coating is an effective barrier or not against the penetration of contaminants. The coating with the highest Ge additions was too brittle to survive the 303 — a— Si-Ge/A1F3/A1203 • • e - - Si-Ge/MgF2/SiC — • • • - Uncoated CP Ti - *- - Si-Ge/CuF2/A1203 Uncoated CP Ti AlfyActivated 5 1.5- i h 6 u § ja - e — !1 U 4 - > _ o - -e- — O — n / CuF-Activated . a> I / “ ♦ " .^1 .-.vO5*' MgF -Activated I 1 I 1 I ’ ’ I 1 1 1 I 1 I T 40 80 120 160 Number of Cycles Figure 128. Cyclic oxidation at 800°C of three different Ge-doped silicide coatings grown on CP-titanium at 1150°C for 12 hours by packs comprised of: (1) Si-Ge/MgF2/SiC, (2) Si-Ge/A1F3/Al203, and (3) Si-Ge/CuF2/Al203. 304 Table 63. A summary for the detection of oxygen penetration by microhardness and etching for Ge-doped silicide coatings grown on CP-titanium by three different packs at 1150°c for 12 hours following cyclic oxidation at 500°C, 600#C, 700°C, 800°C, 900°C and 1000°C. The number in brackets is the number of coupons tested: (1) Si-Ge/MgFj/SiC (2) Si-Ge/A1F3/Al203 (3) Si-Ge/CUF2/Al203 Key: EB - Coating an effective barrier, substrate contamination was not detected. PF » Partial failure of the coating at 1-5 regions with observable substrate contamination. TF ■> Total failure of coating was substrate contamination. Temp MgF2 a i f 3 c u f 2 1000°C TF [4] TF [4] TF [4] 900°C TF [4] TF [4] TF [2] PF [3] 800°C TF [4] EB [5] EB [4] 700°C PF [4] EB [1] EB [5] EB [4] 600°C EB [4] EB [5] EB [4] 500°C EB [4] EB [5] EB [4] 305 thermal cycling at high temperature. Additionally, the coatings formed with the A1F3 and CuF2 activators contain large amounts of condensed activator at the coating surface, as may be required for good cyclic oxidation resistance. The shape of the cyclic oxidation curves for the Ge- doped silicide coatings formed by A1F3- and CuF2-activated packs are approximately the same, but the weight gains are lower for the coatings formed with the CuF2 activator. The coating formed by the CuF2 activator contained more Ge, which indicates that the Ge additions are required. The coatings formed with the cuF2-activated pack contained less convolutions than those formed with the A1F3 activator. Therefore, the improved coating morphology for the silicide coatings formed by a CuF2-activated pack probably produced the slower cyclic oxidation kinetics. The wide variation in cyclic oxidation kinetics for coatings formed by CuF2 and A1F3 activators result from the convolutions in these coatings. However, these coatings prevented oxygen dissolution for cyclic oxidation at 800°C. The cyclic oxidation kinetics at 500°C, 600°C, 700°C, 900°C and 1000°C for the Ge-doped silicide coatings formed by MgF2- , Cu F2-, and AlF3-activated packs are shown in Fig. 173 in Appendix E, and the general trends are discussed in the following section. At 700-1000°c the cyclic oxidation kinetics for CP-titanium coated with the Ge-doped silicide were slower than for the uncoated substrates. However, the cyclic oxidation kinetics for uncoated CP-titanium were slower at 500°C and 600°C, in agreement with the resultB for the B-doped silicide coating. The cyclic oxidation kinetics for the Ge-doped silicide coatings grown by a MgF2-activated pack were slower than the coatings formed by A1F3 and CuF2 activators at 500-800°C, but were faster at 900°C and 1000°C. However, the Ge-doped silicide coatings formed by the MgF2-activated pack were only partially protective at 700°C , and totally protective at 600°C and 500°C. The Ge- doped silicide coatings formed by the A1F3- and CuF2- activated packs were protective at 500-800°C, and some were partially protective at 900°C. Since 900°C is above the a to 0 allotropic transformation temperature for titanium (882°C), the stresses associated with this transformation may have produced the failures. In order to separate the transformation stresses from the thermal stresses, the Ge- doped silicide coatings were tested by cyclic oxidation at 875°C, as shown in Fig. 129. The cyclic oxidation kinetics for the Ge-doped silicide coatings formed by a CuF2- activated pack were approximately parabolic, but the kinetics for the AlF3-activated pack showed an increasing rate after 150-200 cycles. The Ge-doped silicide coatings formed with the CuF2 activator were protective, but substrate contamination was observed for the coatings formed with the AlF3-activator, The Ge-doped silicide coatings 307 grown by a CuF2-activated pack were not so convoluted which explains the improved protection by these coatings. The amount of Ge added to the silicide coating is controlled by the ratio of Si to Ge used in the powder pack to form the coating, and the activator stability, as shown in Figs. 68-70. Since the morphology of the coating is affected by changing the halide activator and filler used to form the coating,. altering the Si to Ge ratio for each pack composition gives an unambiguous method for studying the * effect of the Ge additions on the cyclic oxidation kinetics. The cyclic oxidation kinetics at 800°C for Ge-doped silicide coatings with Si to Ge ratios of 2:1, 1:1 and 1:2 that were grown by MgF2-, A1F3-, and CuF2-activated packs at 1150°C for 12 hours are summarized in Figs. 130, 131 and 132. The slowest cyclic oxidation kinetics were always observed for the Ge-doped silicide coatings formed with a Si to Ge ratio of 2:1, which also contained the least amount of Ge. The cyclic oxidation kinetics at 500°C, 600°C, 700°C, 900°C and 1000°C for the Ge-doped silicide coatings formed with Si to Ge ratios of 1:2, 1:1, and 2:1 by MgF2-, A1F3-, and CuF2- activated packs are compared in Figs. 174, 175, and 176 of Appendix E. The cyclic oxidation kinetics for the Ge-doped silicide coatings formed with a Si to Ge ratio of 2:1 were generally the slowest for each activator at every temperature studied. The Ge-doped silicide coatings formed with higher Si to Ge ratios in the pack, which also 308 contained more Ge, were more convoluted. In fact, convolutions were observed for coatings grown by a pack composed of (1:2) Si-Ge/MgF2/SiC, but no convolutions resulted for the (2:1) Si to Ge ratio of the same pack powder. Thus, higher Ge additions produced a larger change in molar volume upon diffusive conversion of the substrate into the Ge-doped silicide coating, which produced more convolutions. The convolutions are easy sites for the nucleation and growth of cracks, and faster cyclic oxidation kinetics. A summary for the detection of substrate contamination for the Ge-doped silicide coatings with different Si to Ge ratios is given in Table 64. The Ge- doped silicide coatings with lower si to Ge ratios were less protective, and the coating failures and substrate contamination were usually observed at the convoluted regions. Therefore, smaller amounts of Ge additions are desirable for the best cyclic oxidation resistance. The cyclic oxidation kinetics for the undoped silicide coatings are also compared with the Ge-doped silicide coatings in Figs. 130, 131 and 132, and in Figs. 177, 178 and 179 of Appendix G. The cyclic oxidation kinetics for the undoped silicide coatings are generally faster than the Ge-doped silicide coatings with the smallest amount of Ge additions (2:1 Si to Ge ratio) at every temperatures examined in this study. Comparison of the results for the detection of substrate contamination (Tables 60 and 63) 309 0.8 •; He •§> £ u 0.6 •: A1F-Activated Pack QO § 6 0.4 H § I 0.2 CuF„-Activated Pack 0 40 80 120 160 200 Number of Cycles Figure 129. Cyclic oxidation at 875°C of two different Ge- doped silicide coatings grown on CP-titanium at 1150°C for 12 hours by packs comprised of: (l) Si-Ge/A1F3/Al203, and (2) Si-Ge/CUF2/Al203. 12Si-6Ge/MeF2/SlC Uncoated CP Ti - - 12SI-12Ge/M*F2/SiC ■ - • -* •6Sl-l2Ge/M;• 6SM2Ge/MsP2/SiC 1.2 — X- • • Sl/MgF2/SiC Uncoated CP TI 1I S' f<-l:2 Ge-doped 'r' 1 1:1 Ge-dope<£ 0.8 .1 Unddpe 1 Silicide & i je-doped ^ V " V w ' ^ 2:1 0.2-i I I I | i ;- t I I 1 ’ 1 1 I 1 I 40 80 120 160 Number of Cycles Figure 130. Cyclic oxidation at 800°C of Ge-doped silicide coatings grown by a MgF,-activated pack on CP-titanium at 1150°C for 12 hours with three different Si to Ge ratios in the powder pack; (2:1), (1:1) and (1:2). Also shown is undoped silicide coating and uncoated CP-titanium. — -— 16Si*8Ge/AIF3 • • o - . Uncoated CP Ti - - * 16Si*16Ge/AlF3 * * • - * 8Si-16Ge/AIF3 —x *Si/AIF3 Uncoated CP Ti 2:1 Oe*doi 1:1 Ge-doppd-* ~~ * 2:1 Gc-doi "1.'fOe-doped ‘~~~Undo^d ,S Uicido< N u m b e r o f C y c le s Figure 131. Cyclic oxidation at 800°C of Undoped and Ge- doped silicide coatings grown by an AlF3-activated pack on CP-titanium at 1150°C for 12 hours with three different Si to Ge ratios in the powder pack; (2:1), (1:1) and (1:2). — 16Sl>8Ge/CuF2 • • • • • Uncoiled CP TI - - • 16SM60e/CuF2 8Si.J60e/CuF2 —x-..Si/CuP2 > Undoped SUlcidc; Uncoiled CP Ti ^ Ah « t 6 X 1:1 Gc-dopgd*. -t» - 2 je-doped Z / t ;' 1 t ■ * *.v 2:1 Ge-dope«i' ' ZZ, ---- *" mTje-3ow 1 1 I 1 1 ’ I ’ I » I ■ I ' l ■ | ■ | >■ 0 40 80 120 160 200 Number of Cycles Figure 132. Cyclic oxidation at 800°C of Undoped and Ge- doped silicide coatings grown by a CuF2-activated pack on CP-titanium at 1150°C for 12 hours with three different Si to Ge ratios in the powder pack; (2:1), (1:1) and (1:2). 311 Table 64. A summary for the detection of oxygen penetration by microhardness and etching for Ge-doped silicide coatingB grown on CP-titanium by three different packs with two ratios of Si to Ge at 1150°C for 12 hours following cyclic oxidation at 500°C, 600°c, 700°C, 800°C, 900°C and 1000°C. Two coupons were tested for each pack: (1) 12 Wt.% Si, 12% Ge, 2% MgF2, SiC [1:1] (2) 6% Si, 12% Ge, 2% MgF2, sic [1:2] (3) 16% Si, 16% Ge, 2% A1F3, A1203 [1:1] (4) 8% Si, 16% Ge, 2% A1F3, A1203 [1:2] (5) 16% Si, 16% Ge, 2% CUF2, A1203 [1:1] (6) 8% Si, 16% Ge, 2% CUF2, A1203 [1:2] Key: EB ° Coating an effective barrier, substrate contamination was not detected. PF Partial failure of the coating at 1-5 regions with observable substrate contamination. TF » Total failure of coating was substrate contamination. Temp [1:1] MgF2 [1*2] MgF2 [l:l] 1000®C TF TFTF 900°C TF TF TF 800°C TF TF EB 700°C TF TFEB 600°C PF TF PF 500°C EB EB EB [1:2] A1F3 [1:1] c u f 2 [1:2] 1000°C TF TFTF 900°C TF TFTF 800°C TF EB TF 700°C TF EBPF 600°C PF EB PF 500°C EB EB EB 312 demonstrates that Ge-doped silicide coatings are more * protective at 600°C to 800°C. Although the Ge-doped silicide coatings contained some convolutions, the Ge additions are required to provide cyclic oxidation resistance at high temperature. Thinner Ge-doped silicide coatings were grown by processing at 950°C for 6h and 28h, and the cyclic oxidation kinetics at 700°C and 800°C are compared with undoped and the thicker coatings formed by the MgF2, A1F3 and cuF2 activators in Figs. 133, 134, and 135, respectively. The cyclic oxidation kinetics at 500°C, 600°C, and 900#C for the thinnest Ge-doped silicide coatings are compared with the thick Ge-doped silicide coatings formed by MgF2, A1F3, and Cu F2 activators in Figs. 177-179 of Appendix E. The cyclic oxidation kinetics for the thinnest Ge-doped silicide coatings are generally faster than the thicker coatings at 700-900°C. Additionally, the thinnest Ge-doped silicide coating formed at 950°c for 12 hours by A1F3- and CuF2- activated packs were highly convoluted and not very adherent. The cyclic oxidation resistance for these thin Ge-doped silicide coatings was worse than the thick coatings at every temperature. The summary for the detection of substrate contamination in Table 65 shows that the thinner coatings were less protective than the thicker coatings at all temperatures. The cyclic oxidation kinetics for the slightly thicker Ge doped silicide coatings, which are 313 Uncoated CP Ti "6 (20-15 tim) & E o bo (44-35 |im) (20-15 um) JZ BO 5 (90-80 um) »~r ~» i 1 n I 1 I * I 1 \ + 40 80 120 160 200 (a) Number of Cycles Uncoated CP Ti "e (44-35 um) •& BOt> (20-15 um) ”5) I (90-80 um) 80 120 160 (b) Number of Cycles Figure 133. Cyclic oxidation of the undoped and Ge-doped silicide coatings grown on CP-titanium by a MgF2-activated pack at 1 1 5 0 ° C for 1 2 hours, 9 5 0 ° c for 2 8 hours, and 9 5 0 ° c for 6 hours tested at (a) 7 0 0 ° C and (b) 8 0 0 #C for 2 0 0 cycles. 314 5*j e (15-11 um) ■5 4*1 E ti 00 / 3-1 * - *•*' 2 \ /S' Uncoati ;d CP Ti s (40-32 |lm) . • o • •* I (a) 40 80 120 160 200 Number of Cycles "e (15-11 Jim) 1 o Uncoated CP Ti 00 : (80-60 Jim) (40-32 iim) i 'nr ■ i ■ r » i » i ■ i i i i r 40 80 120 160 (b) Number of Cycles Figure 134. Cyclic oxidation of the undoped and Ge-doped silicide coatings grown on CP-titanium formed by an A1F3- activated pack at 1150°C for 12 hours, 950°C for 28 hours, and 950°C for 6 hours tested at (a) 700°C and (b) 800°C for 200 cycles. 315 "E (25-20 Jim) ■§» E, Uncoatec CPTi oDO -C OO 5 (75-65 uin) 40 80 120 160 200 (a) Number of Cycles 6-7 (25-20 Jim) "E 4.8-1 ■§> wE U 3.6 \ - Undated CP Tis 0 0 2.4-! / s' (46-38 Jim). (Zs3o im) *§) I 1.2 (75-65 Jim) b ^ = T * “ ’ r i ■ i i i i 1 i ■ i ■ i ■ i ■ 40 80 120 160 200 (b) Number of Cycles Figure 135. Cyclic oxidation of the undoped and Ge-doped silicide coatings grown on CP-titanium formed by a CuF2- activated pack at 1150°c for 12 hours, 950®C for 28 hours, and 950°C for 6 hours tested at (a) 700°C and (b) 800°C for 200 cycles. 316 Table 65. A summary for the detection of oxygen penetration by microhardness and etching for Ge-doped silicide coatings grown on CP-titanium at 9S0°C for either 28 hours, 12 hours or 6 hours by three different packs following cyclic oxidation at 500°C, 600°c, 700°c, 800°c, 900°c and iooo°c. Two coupons were tested for each pack: (1) Si-Ge/MgF2/SiC, 950°C, 28 hours (2) Si-Ge/MgF2/SiC, 950°C, 6 hours (3) Si-Ge/A1F3/Al203 950°C, 28 hours (4) Si-Ge/A1F3/Al203 950°C, 12 hours (5) Si-Ge/CUF2/Al203 950°C, 28 hours (6) Si-Ge/CuF2/Al203 950°C, 12 hours Key: GB «* Coating an effective barrier, substrate contamination was not detected. PF ■ Partial failure of the coating at 1-5 regions with observable substrate contamination. TF ■ Total failure of coating was substrate contamination. Temp MgF2, 6h A1F3, 12h CUF; 900°C TF TF TF 800°C PFTF PF 700°C TF TF PF 600°C PF TFTF 500°C EBTF TF NgF2, 28h A1F3, 28h CuF. 800°C PFPF TF 700°C PF PFPF 317 formed at 950°C for 28 hours, are similar to the thicker coatings at 700°C, and slightly faster at 800aC. However, none of the thinner coatings prevented substrate contamination for cyclic oxidation at 700°C and 800°C. Therefore, thicker Ge-doped silicide coatings are required to prevent the penetration of oxygen during cyclic oxidation. A comparison between the cyclic oxidation kinetics for the B- and Ge-doped silicide coatings on CP-titanium (Figs. 113 and 128) demonstrates that the weight gains are about the same. However, the Ge-doped silicide coatings are more protective at temperatures above 800°C, which indicates that a Ge-doped silica scale is more protective at higher temperatures. This observation may be related to the evaporation of boron that is localized at the surface during high-temperature oxidation. However, the thinner B-doped silicide coatings were much more protective than the Ge- doped silicide coatings. Convolutions were not observed for the B-doped silicide coatings, and these coatings may be more resistant to mechanical stresses. In general, the differences between the B- and Ge-doped silicide coatings are not very significant. A comparison for the CTE versus temperature for Tisi2 [163], titanium [164] and Ti-24Al-llNb [14] is given Fig. 136. The CTE for Tisi2 is almost exactly the same as Ti up to the allotropic transformation temperature for Ti. Since 318 TiSi2 Is a brittle intermetallic compound, large stresses # resulting from a CTE mismatch could not be tolerated. However, the CTE match for TiSi2, which is a major constituent for the titanium-silicide coatings, with Ti is one reason for the excellent cyclic oxidation resistance of these coatings. The similarity in CTE for TiSi2 and the de base Ti-24Al-llNb is one reason for the excellent cyclic oxidation resistance of the coatings on these alloys. 5.3.4 Geruaniun-Doped Silioida coatings on Ti-Al-Nb Alloys The Ge-content for the Ge-doped silicide coating is varied by using a MgF2, A1F3, or CuF2 activator to form a thick coating at 1150®C for 12 hours on Ti-22Al-27Nb and Ti- 20Al-22Nb (Figs. 74 and 76). The cyclic oxidation kinetics at 800°C in Fig. 137 for the Ge-doped silicide coatings grown by an AlF3-activated pack on Ti-22Al-27Nb and T1-20A1- 22Nb are the same. Thus, the substrate has little influence on the performance of the Ge-doped silicide coating for the orthorhombic alloys, in agreement with the B-doped silicide coatings. The cyclic oxidation kinetics for the Ge-doped silicide coatings are approximately parabolic. However, larger weight gains were observed for the Ge-doped silicide coatings formed by a CuF2 activator. The cyclic oxidation kinetics at 500°C, 600°C, 700°C, 900°C and 1000°C are given in Fig. 180 of Appendix E, and the results are summarized in the following discussion. The cyclic oxidation kinetics for the Ge-doped silicide coatings 319 o• • • OT1-24AI—11Nb □— □Titanium o— oTiSI, p 20,0 - c o '55 c o 13.0- Ti—24AI— 11Nb, -° a . l2 a o- * ‘ * T1SI2 g i o " « TltoniUm . -o' 10.0- i£ o r e-*- c V "o £I) 5.0------1------\------1------1------1------1------1-----—i------0 100 200 300 400 500 600 700 600 900 1000 s Temperature (C) Figure 136. Coefficient of thermal expansion versus temperature for Ti [167], TiSi2 [166] and Ti-24Al-llNb [10]. H r* ^ 7 —Uncoated Ti-20Al-22Nb _ _ "6 Uncoated Ti-22AI-27Nb 0.8 -j ■&> r t ! B, , CuF,-Activated (20-22; o oo 0.6 ■ ; 5 6 I// / 3 : M AIF-Activated (20-22) AlFj-Activatcd (22-27) £ -I I gF,-Activated (22-27) 80 120 160 Number of Cycles Figure 137. The cyclic oxidation kinetics at 800°C for the Ge-doped silicide coating grown on Ti-22Al-27Nb and Ti-20A1- 22Nb by a MgF2-, A1F3- and CuF,-activated pack at 1150°C for 12 hours. 320 were slower than for the uncoated alloys at 700-1000°C. However, the cyclic oxidation kinetics for some of the Ge- doped silicide coatings were faster than the uncoated alloys at 600°C and 500°C, in agreement with the results observed for the B-doped silicide coatings. The thick Ge-doped silicide coatings had convolutions at the corners, and were susceptible to spalling by the low-temperature oxidation mechanism described for the B-doped silicide coatings (section 5.3.2). The spalling is shown by the large weight losses for low-temperature cyclic oxidation. However, the Ge-doped silicide coatings formed by the AlF3-activated pack were resistant to spalling during low-temperature cyclic oxidation. A summary for the detection of oxygen contamination by microhardness and etching is given in Table 66. The Ge- doped silicide coatings were protective at every flat area of the substrate for every temperature used in this study. However, failure was observed for low-temperature oxidation at the convoluted corners. The Ge-doped silicide coatings formed by MgF2 and CuF2 activators had the most convolutions at the corners, and the most failures were observed for these Ge-doped silicide coatings. Additionally, the Ge- doped silicide coating grown by an AlF3-activated pack on Ti-20Al-22Nb was susceptible to corner attack. Thin Ge-doped silicide coatings were formed at 950°C for 6 hours or 28 hours by either a MgF2-# A1F3-, or 321 Table 66. A summary for the detection of oxygen penetration by microhardness and etching for Ge-doped silicide coatings grown on Ti-22Al-27Nb (22-27) and Ti- 20A1—22Nb (20-22) at either 1150°C for 12 hours, 950°C for either 28 hours or 6 hours by three different packs following cyclic oxidation at 500°C, 600°C, 700°C, 800°C, 900ttC and 1000°C. (1) Si-Ge/MgF2/SiC, on Ti-22Al-27Nb, 1150°C, 12 hours (2) Si-Ge/MgF2/SiC, on Ti-20Al-22Nb, 950°C, 28 hours (3) Si-Ge/MgF2/SiC, on Ti-20Al-22Nb, 950°C, 6 hours (4) Si-Ge/A1F3/Al203 on Ti-22Al-27Nb, 1150°C, 12 hours (5) Si-Ge/A1F3/Al203 on Ti-20Al-22Nb, 1150°C, 12 hours (6) Si-Ge/A1F3/Al203 on Ti-20Al-22Nb, 950°C, 28 hours (7) Si-Ge/A1F3/Al203 on Ti-20Al-22Nb, 950°C, 6 hours (8) Si-Ge/CuF2/Al203 on Ti-20Al-22Nb, 1150°C, 12 hours (9) Si-Ge/CuF2/Al203 on Ti-20Al-22Nb, 950°C, 28 hours (10) Si-Ge/CuF2/Al203 on Ti-20Al-22Nb, 950®C, 6 hours Key: EB *» Coating an effective barrier, substrate contamination was not detected. PF - Partial failure of the coating at 1-5 regions with observable substrate contamination. TF ■ Total failure of coating was substrate contamination. Temp MgF2, a i f 3, CuF2 (1) (4) (5) (8) 1000°C EB EB EB PF 900°C EB EB EB EB 800°C EB EB PF PF 700°C EB EB EB PF 600°C PF EB EB PF 500°C PF PF EB PF MgF2, a i f 3, C u F2 (3) (?) (10) 1000°C TF TF TF 900°C PF EB EB 800°C EB EB EB 700°C EB EB EB 600°C EB EB EB 500°C PF EB PF MgF2, a i f 3, C u F2 (2) (6) (9) 1000*C EB EB TF 800°C EB EB EB 500°C PF EB PF CuF2-activated pack, and the cyclic oxidation kinetics at 1000°C and 800°C are compared with the thicker coatings in Figs. 138, 139, and 140, respectively. The cyclic oxidation kinetics at 800°C are generally the same for all of the Ge- doped silicide coatings. However, fast, linear cyclic oxidation kinetics were observed at 1000°C for all of the thinnest Ge-doped silicide coatings. The summary of oxygen contamination detection for the thin Ge-doped silicide coatings given in Table 66 indicates that oxygen contamination and coating failure were observed for all of the thinnest Ge-doped silicide coatings. The thin Ge-doped silicide coatings formed at 950°C for 28 hours by MgF2 and A1F3 activators were protective at 1000°C and slow cyclic oxidation kinetics were observed. However, coating failure and a fast rate of cyclic oxidation are observed for the Ge- doped silicide coating formed by a CuF2-activated pack. The Ge-doped silicide coatings grown by a CuF2 activator at 950°C (Figs. 78-80) had the highest Ge additions, and the coating was more brittle and slightly convoluted at the corners. Thus, large amounts of Ge additions are not desirable for the Ge-doped silicide coatings, in agreement with the results observed for the Ge-doped silicide coatings on CP-titanium. The cyclic oxidation kinetics at 1000°C for a thin Ge- doped silicide coating grown on a Ti-24Al-llNb alloy by a AlF3-activated pack are shown in Fig. 139. Rapid oxidation 323 " 6 . Uncoated Ti-22Al*27Nb *§> £ o (20-22:24-18 Jim) / 00 (20-22:40-34 Jim) § I Uncoated Ti-20A1-22Nb &2*27: *50-120Jim) 1~ri ip i'i | i i i | i i i | ii I | i i i |i ii | i i i 40 80 120 160 200 (a) Number of Cycles ■ Uncoated Ti-20Al-22Nb 0.35- Uncoated Ti-22Al-27Nb "S ■§> 0.25- (20-22:24-18 um) (20-22:40-34 Jim) § 6 0.15- I) £ (22-27:150-120 um) i i i i » i » ~r * i 1 i 1 i 1 i 1 r 0 40 80 120 160 200 (b) Number of Cycles Figure 138. Cyclic oxidation of Ge-doped silicide coatings grown on Ti-22Al-27Nb and Ti-20Al-22Nb by a MgF2-activated pack at either 1150°C for 12 hours, 950°c for 28 hours, or 9S0°C for 6 hours, and tested at (a) 1000°C and (b) 800°c for 200 cycles. 324 1i Uncoated Ti-22Al*27Nb 2- i WE i 1.6- (20*22:39-33 Jim) />' *§> i E * 4> (24-11:25-15 Jim) P (20-22:22-1 6 0 1.2- § t 6 0.8- f /1 i\ V k ' _ V - \ ^ '2 -2 7':100,80 ' m) § it I 0.4- i - \(20p22:80^0 Jim) • *^- Uncoated Ti-20Al-22Nb 0 .35- ■Uncoated Ti-22Al*27Nb *¥ (20-22: 80-60 Jim) ■§» (20-22:39-33 Jim) j= 0 .25- 41 60 § 0.2*: e v / - 0 .1 5 -iJ I (22-27: 100-80 Jim) (22-27:22-18 Jim) I 1 I 1 I "l~"l 1 I 1 I 1 I 1 I 40 80 120 160 200 (b) Number of Cycles Figure 139. Cyclic oxidation of Ge-doped silicide coatings grown on Ti-22Al-27Nb and Ti-20Al-22Nb by an AlF3-activated pack at either 1150°C for 12 hours, 950°c for 28 hours, or 950°c for 6 hours, and tested at (a) 1000°C and (b) 800°c for 200 cycles. 325 5 4 B Uncoated Ti-22Al-27Nb ! 3 (20-22:13-9 jitn) ei) 3 6 2 4-4 § (20-22:33-27 fitn) I 1 = t * * (2 0 -2 2 : 100-80 n m ) 0 Uncoated Ti-20Al-22Nb 40 80 120 160 200 (a) Number of Cycles U n co ated H -2 0 A l-2 2 N b Is Uncoated Ti-22Al-27Nb S. fli t>0 — ■ (20-22MQ0-8Q § 8 s J r !*__ _♦ / ■ • 1) i “ - : V ^^(20-22:13-9 Mm) i 0.2 *= : t o r n ------* ■ - « ------i f 0' (20-22:33*27 pm) ° ' T 1 I 1 I 1 I "1 T 1 I r -\ i | i | i | i 0 40 80 120 160 200 (b) Number of Cycles Figure 140. Cyclic oxidation of Ge-doped silicide coatings grown on Ti-22Al-27Nb and Ti-20Al-22Nb by a CuF2-activated pack at either 1150°C for 12 hours, 950°C for 28 hours, or 950°C for 6 hours, and tested at (a) 1000°C and (b) 800°C for 200 cycles. 326 kinetics and substrate contamination were observed for these substrates, in agreement with the results for the Ge-doped silicide coatings on the Ti-20Al-22Nb alloy. The thin Ge-doped silicide coatings formed at 950°C for 28 hours provide high-temperature cyclic oxidation resistance that is similar to the thick coatings, but the low-temperature cyclic oxidation resistance must be examined. The cyclic oxidation kinetics at 500°C for the thick and thin Ge-doped silicide coatings formed by MgF2-, AIF3-, and CuF2-activated packs are shown in Figs. 141, 142, and 143, respectively. The large weight losses show the spalling of the thick Ge-doped silicide coatings. No spelling was observed for the thinner Ge-doped silicide coatings, and only small weight gains were observed. However, the Ge-doped Bilicide coating formed by a CuF2- activated pack at 950°C for 28 hours did spall at the convoluted corners, as shown by the weight loss in Fig. 143. The cyclic oxidation results at 600°C, 700°C and 900°C for the thinner Ge-doped silicide coatings formed by MgF2-, AIF3- and CuF2-activated packs are given in Appendix E (Figs. 181, 182, and 183, respectively), and are summarized in the proceeding discussion. The thin Ge-doped silicide coatings were protective and resisted spalling during low- temperature cyclic oxidation. However, the thin Ge-doped silicide coatings were not are protective in cyclic oxidation at 900°C and 1000°C. The thin Ge-doped silicide 327 TT*\ (22-27:150-120 Jim) (20-22:24-18 Jim) Uncoated T1.22AI-27Nb M -6- Uncoated T1-20AI-2: :Nb 8 - -12 0 40 80 120 160 200 Number of Cycles Figure 141. Cyclic oxidation at 500°C for Ge-doped silicide coating grown on Ti-22Al-27Nb and Ti-20Al-22Nb by a MgF2- activated pack at 1150°C for 12 hours, 950 °C for 28 hours or 6 hours. 0.2 _ _ o y fiffv f t - >?j ■5"^ — \ “ "4 - __ i I \ -0.2 ' ' " “ ' f UQCoatedTl.204t.22Nb ■ (22-27: 100-80 Jim) eb -0.4 E (20-22:39-33 Jim) u -0.6^ Uncoated Tl-22Al-27Nb a% - 0.8 5 -1-2^ (20-22:80-60 Jim) I -1.4? -1.6 ? - 1.8 4 ■ | < Figure 142. Cyclic oxidation at 500#C for Ge-doped silicide coating grown on Ti-22Al-27Nb and Ti-20Al-22Nb by an A1F3- activated pack at 1150°C for 12 hours, 950°C for 28 hours or 6 hours. 328 (20*22:33*27 Jim), (20*22 B Uncoated Ti-20/ l* 22Nb Uncoated Ti*22Al*27Nb & (20s22:100-80 Jim) & s ■§> 'ri'i'i ri'i |111'i | 11 i11 i i 11 i ii 11 i i | iti'i i ii | r r r 80 120 Number of Cycles Figure 143. Cyclic oxidation at 500°C for Ge-doped silicide coating grown on Ti-22Al-27Nb and Ti-20Al-22Nb by an CuF2- activated pack at 1150°C for 12 hours, 950°C for 28 hours or 6 hours. ERNEST F. FULUtM, INC. nn m tat Figure 144. Optical photograph of B- and Ge-doped silicide coatings grown on Ti-22Al-23Nb fatigue specimens at 950*0 for 6 hours, which are marked as "a" and "b", respectively. coatings formed by A1F3- and CuF2-activated were the only coatings able to protect the alloy substrates during cyclic oxidation at 900°C. CHAPTER VI SUMMARY AND CONCLUSIONS 1. Titanium-silicon diffusion couples were used to determine the interdiffusion coefficients for TiSi2, TiSi, Ti5Si4, TigSij, and Ti3Si. Equations were derived from the theory of Wang et al. [112] to establish the intrinsic parabolic rate constants for five-layered growth. Comparison of the measured data for TiSi2 and TiSi revealed that the growth is controlled by solid-state diffusion in the bulk. 2. The growth kinetics for the five-layered TiSi2/Tisi/TisSi4/TisSi3/Ti3Si diffusion coatings by fluoride-activated packs were compared with the rates measured by the diffusion couple. The rate-limiting step for the growth of the undoped five-layered coating by the less stable A1F3 and CuF2 activators is solid-state diffusion. Lower Si-fluoride vapor pressures were produced by the more stable MgF2 activator, and the rate-limiting step became gas-phase diffusion. Boron was not dissolved into the silicide coating, and the growth kinetics for the B-doped silicide coating were similar to the undoped silicide 330 331 coatings. Germanium was dissolved into the silicide coating # layers, and the growth kinetics for the Ge-doped silicide coatings were different than for the undoped silicide coatings. 3. The kinetics for the growth of dual-layer TiB2/TiB boride diffusion coatings from A1F3- and NgF2-activated packs were measured. The thickness of the inner TiB layer is nonuniform, and does not follow a parabolic growth law. The rate-limiting step for the growth of the TiB2 layer was solid-state diffusion. 4. B-doped silicide coatings were developed for CP-titanium, Ti-22Al-27Nb, Ti-20Al-22Nb and Ti-24Al-llNb by a fluoride- activated, pack cementation method. A boride compound (TiB2) was used to lower the B activity in the MgF2- activated pack, which produced silicide coatings with a TiB2 layer at the outer surface. The thickness of the TiB2 layer was increased by: (1) increasing the B activity of the boride compound (TaB2, CrB2) in the pack, or (2) using a less stable activator (A1F3, CuF2). 332 5. Ge-doped silicide coatings were developed for CP- titanium, Ti-22Al-27Nb, Ti-20Al-22Nb and Ti-24Al-llHb by three different packs: (1) 12 wt.% Si, 6% Ge, 2% MgF2 and sic, (2) 16% Si, 8% Ge, 2% A1F3 and A1203 and (2) 16% Si, 8% Ge, 2% CuF2 and A1203. The Ge content in the coating was varied by the ratio Si to Ge, choice of halide activator, and the inert filler of the powder pack. 6. The isothermal oxidation kinetics of the Ge-doped and B- doped silicide coatings were always slower than pure TiSi2, and equivalent to or faster than Si02. The activation energy for oxidation at 700-1000°C was generally equivalent to the growth of Si02 for the B-doped and Ge-doped silicide coatings. The isothermal oxidation kinetics for the Ge- doped silicide coatings were slower than for B-doped silicide coatings. The activation energy for the low- temperature oxidation (700-500°C) was low, which indicates that solid-state diffusion along Ti02 grain boundaries that are doped with impurities was the rate-limiting step. Analysis of the oxide scales agrees with the kinetic data: (1) a two-layer scale was observed for the high-temperature oxidation, with an inner layer of almost pure Si02 adjacent to TiSi2 and an outer layer of Ti02-Si02, (2) a single-layer scale of mixed Ti02-Si02 was observed following low- temperature oxidation. 333 7. Cyclic oxidation for the B- and Ge-doped silicide coatings on CP-titanium showed that the best coating prevented substrate contamination at 500-800°C for 200 oxidation cycles. The best B-doped silicide coatings were formed by packs composed of either Si-TiB2/MgF2/Al203 or Si- TiB2/CuF2/Al203 at 1150°c for 12 hours, which had the lowest amounts of TiB2 in the coating. The best Ge-doped silicide coatings were formed by A1F3- and CuF2-activated packs at 1150°C for 12 hours, which contained lower amounts of Ge. 8. Cyclic oxidation of the B- and Ge-doped silicide coatings on Ti-22Al-27Nb and Ti-20Al-22Nb showed that the best coatings prevented substrate contamination at 500-1000°C for 200 oxidation cycles. The best B-doped silicide coatings were grown by packs composed of either Si-TiB2/MgF2/Al203 or Si-TiB2/CuF2/Al203 at 950°C for 28 hours, which had lower amounts of TiB2 in the coating. The best Ge-doped silicide coatings were formed by an AlF3-activated pack at 950°C for 28 hours, which contained a moderate amount of Ge. 9. The growth of undoped and Ge-doped MoSi2 diffusion coatings by the HAPC method results in a byproduct layer of salt-oxide at the surface, whose composition depends on the specific activator used to form the coating. MoSi2 coatings that were grown by the NaF-activated pack did not pest after 2500 hours of isothermal oxidation or 600 1-hr. oxidation cycles at 500°C. The formation of a fast growing Na- silicate scale passivates MoSi2 to prevent accelerated oxidation at low temperature. NaF, NaCl, NaBr, and Nal were applied to bulk MoSi2 and MoSi2 diffusion coatings by an aqueous solution and they also prevented the accelerated low-temperature oxidation of these substrates which otherwise were prone to posting. CHAPTER VIZ FUTURE WORK The excellent oxidation resistance of the B- and Ge- doped silicide coatings has been shown. However, a coated part must also endure some form of mechanical loading in a real engineering application. Since Ti-sillcide coatings are brittle, cracking and fatigue under a load are important issues that must be resolved by mechanical testing of coated specimens. Figure 144 shows Ti-22Al-23Nb fatigue specimens coated with a B- and Ge-doped silicide coatings that will be tested at NASA Lewis Research Center (W. Brindley) in fatigue at a low temperature. If the coating does not reduce the fatigue life of the substrate, then further thermomechanical fatigue testing is required to demonstrate the potential of these coatings. 335 APPENDIX A X-RAY DIFFRACTION OF AS-COATED WORKPIECES III •0 to •0■0 - 90M mtctw TrtT W w fl«tw f CWtfOw, (a) »T" IJ4* to - 90 • 10 TITmll* OMCIM troll t r w (b) Figure 145. Undoped silicide coatings grown on CP-titanium at 1150°C for 12 hours by packs composed of (a) 10 wt,% si, 2% A1F3 and A1203, and (b) 10% Si, 2% MgF2 and A1203. 336 337 4.410 a.y?0 a.ysa i.pa 1.941 l i l i . t.fM* • 00 174.■ eii.7 (0 2 0 ) TiSi 70 ■04.0 ■ 00 170.9 • SO l».4 1 40 102.1 H . a- (004) TISI \-to 14.1' A* ~ * *• *i i .11 ■ 10 T n T i’j “»' i i i i i i i i j i i i."ii i i i I i i i i 1 CO 10 40 00 00 70 00 TITANIUM BlLlCIOE 17*0434 J TI L El TITANIUM SlLtCID! 17*0907 JLL T13 014 titanium irucroc ■1*1070 u . _LL (a) TI0 014 1.^41 (316) TI SI 00 (0 0 4 ) TISI 90 111]) TI 01 K (112) TV SI CO » 90 40 TITANIUM 01L1C1OE TIB 014 TITANIUM 9H.ICIPC >0-1900 TITANIUM ULICIOe (b) TI 01 Figure 146. Undoped silicide coating that was sanded prior to XRD to reveal the inner layers (a) TiSi layer and (b) Ti5sl4 and Ti5si3 with some TiSi. 338 :p s S 941.0' 1 1 100 499.B' aiB,4870.9 loa.a too. a- 94.1' 0.0' i" i i ■ i iti—i i i i 11 i-r i— i t 1 — i— i i i i— i j p i to 90 40 90 90 70 00 TITANIUM StLlCIOE 39*0709 ll 1 ... ______t i sia TITANIUM BOMIDE 39-0741 . 1 ----- 1--- 1------1...... -I. ti oa (a) fcpt «•**• 1.979 l.fM 1.CL i . CL i.CL t.lB9| 199.0'----- 1— ■100•0 90 117.9' 90 101.0'\ . |40 79.9' 90 91.0' 10 19.9' to 0.0 0 10 90 40 ' 'Jo' ' ,m'~3Tm' - wo --- J, TITANIUM HUCIOe 39-0709 ll 1 ...... TI BIO TITANIUM BORIDE 39-0741 . i i 1 1 1 TI Bt TANTALUM OILICIOC 39-0493 . /. . 1 . . ______1______1 . 1 1 1 (b) T* Oil Figure 147. Boron-doped silicide coatings grown on CP- titanium at 1150°C for 12 hours by packs composed of (a) 7 Wt.% Si, 6% TiB2, 2% MgF2 and A1203, (b) 7% Si, 6% TaB2, 2% MgF2 and A1203, and (c) 7% Si, 6% CrB2, 2% MgF2 and A1203. Figure 147. Continued. 8.f78 c 1.741 __4f" p * I.fU 1.741 1. 884.8- * 80 tts.t- • 80 809.0- • 70 178.4- - 80 147.0- • 80 117.8- - 40 aa.a- 88. 8- ■ 80 88.4- . 1 — IHM p A - 10 0.0J •ri"i to 90 40 " i " ' W o ' a10 TITANIUM tOfltOt 19*0741 . 1 rt at t i t a n i u m stLtcioe 19-0789 ll 1 ...... Tl 111 (c) 340 a.yw a.yaa i.yai i.ptt ■too - 90 90 30 SO SO 90 40 TITANIUM OlLICIOe 10*0233 TI BI2 TITANIUM BOniOE TI 92 NIOBIUM SILICIDE (a) BIS 4.430 S. 2.pi i.fn t.ftt t. 474,On !!!!_ ± L -too 429,0 - 90 379.2- * 90 331.8- - 70 ■94.4 • 90 837,0- 90 199 - 40 142 - 20 90 - to S r l J L a J v . 0 " E 30 TITANIUM BILICIOE 10-0229 J l_ j ___ L ii ‘i i TI 912 TITANIUM BDM02 39-0741 TI 92 NIOBIUM SILICIOB 09-0490 1 I L x J u NO SIB Figure 148. B-doped silicide coating formed by a pack composed of 7 wt.% si, 6% TiB2, 2% MgF, and Al203 (a) on Ti- 22Al-27Nb at 1150°C for 12 hours and (b) on Ti-20Al-22Nb at 950°C for 6 hours. 341 4.438 8.978 9.gS2 I.f93 1.^41 1. 9 8 T 389.9 32B.0 987.7 3 4 8 .6 303.9 184.4 193.3 B9.3 t i t a n i u m s i l i c i d e 33-0783 -LJ L. TI 919 TITANIUM SILICIDE 17-0494 J_I II I . 1.1. I 11 I TI 81 TITANIUM SILICIDE 27-0907 ■ I J ■ L J J a-L j . (a) T19 St4 4.438 8.978yte B.yga i.y ?41 1.^43 — > 884.1- I B B , 8 174.3 149.4- 184.0* 99.8 49.0- 80 30 40 TITANIUM SILICIOE 33-0783 II. I TI SIB ALUMINUM OXIDE / COBUNOUM. STM 10-0173 J_L ALB 03 ALUMINUM FLUOR IOE 80-0008 J I LL i .1 I Cb> AL F3 Figure 149. Germanium-doped silicide coatings grown on CP- titanium at 1150°C for 12 hours by packs comprised of (a) 12 wt.% Si, 6% Ge, 2% MgF2 and SiC, and (b) 16% Si, 8% Ge, 2% A1F3 and A1203. 342 -100 - 00 - 00 - TO ' 80 - 80 - 40 * 8030 - 10 80 30 •0 TITANIUM BIUCIOC 10-0888 TI 818 (a) a.yaa t.yai i.f4i |H> **100 00 ■ 80 *• 80TO - 80 • 40 •■ 8030 80 30 •0 TITANIUM 8ILICI0C TI SIS ALUMINUM 0X108 / CORUNDUM. 8VH 03 n Figure 150. Ge-doped silicide coatings grown at 1150°C for 12 hours by a pack comprised of (a) 12 wt.% si, 6% Ge, 2% MgF2 and Sic on Ti-22Al-27Nb, and (b) 16% si, 8% Ge, 2% A1F, and A1203 on Ti-20Al-22Nb. APPENDIX B THICKNESS RATIOS FOR THE GROWTH OF UNDOPED SILICIDE COATINGS 950C 1050C 7 0.5 - a 1I50C CO. Tim e [hours] Figure 151. Plot of the thickness ratio of TiSi2 over TiSi + Ti5Si* + Ti5Si3 + Ti3Si for undoped silicide coatings grown at H50°Cf 1050°C and 950°c by packs comprised of (a) 10 wt.% Si, 2% CuFj and A1203, (b) 10% Si, 2% A1F3 and A1203, and (c) 10% Si, 2% MgF2 and A1203. 343 Figure 151 continued. 151 Figure P' »•=«p/e,^3+54)] p' ip1=(y/G.+vvyi 2 0 0 2 1 3 1 i tm iii (b) (c) 7 1 I 2 2 3 25 20 IS 10 5 1 1 1 1 1 i i 1 1 1 1 1 1 1 1 1 i i 1 1 [ 1 1 u i i i 1 1 1 1 n 1 1 n 1 1 1 n m 1 1 1 n 950C I150C Time [hours] Time Time [hours] Time 1150C 20 050C 050C irri it 344 APPENDIX C PARABOLIC PIT TO THE ISOTHERMAL OXIDATION KINETICS FOR THE BORON- AND GERMANIUM-DOPED SILICIDE COATINGS 0.6 o 2 3 4 5 6 7 Square Root Time [hour]1'1 n m i n u i ■a— 6 0 0 C — e — 6 0 0 C - -G *•1000C • • © • . 700C § 0.0011 — X - - 5 0 0 C □ - A* • 900C -0OOC o --□-•1000C - A- 0.0001 i i i i m i n i i i i i 111| -900C 0.1 1 10 — x ’500C □ 8 0 0 C time [hours] Figure 152. Linear regression fit for the isothermal oxidation data of a B-doped silicide coating grown on CP- titanium at 1150°C for 12 hours by a pack composed of 7 wt.% Si, 6% TiB2, 2% MgF2, and A1203 (a) parabolic plot and (b) log-log plot. 345 346 3 1 0 0 0 C 2.5 rfT' E & 2 E, c *fS 1.5- - O o 0 A .c O o> 1 -I 0 □ 0) : □ 0 .5 -I ° ' b \ ftAffi- n J \ 0^ -rrv 1 j i i i 1 | ■ i i i | t i i i | i i i i | i i i 2 3 4 5 6 (a) Square Root Time [hour]1* 1 0 0 0 C ■§> E, c '5 0 sz ai 1 0.01 i i i i i iii| i iitiiii) 1 10 100 (b) Time [hour] Figure 153. Linear regression fit for the isothermal oxidation data of the B-doped silicide coating grown on CP- titanium at 1150°C for 12 hours by a pack comprised of 7 wt.% si, 6% TaB2, 2% MgF2, and A1203 (a) parabolic plot and (b) log-log plot. 347 900C tr* E ■§, 0.8-i *i 0.6 \ a • 800C 0 .4 -i 600C 0.2 700C 0 1 2 3 4 S 6 7 (a) Square Root Time [hour] 1/2 1 i— i i 11111------1— i— i i i 1111 i i—r i i 111 1 10 100 Time [hour] Figure 154. Linear regression fit for isothermal oxidation data of uncoated Ti-22Al-27Nb (a) a parabolic plot and (b) a log-log plot. 348 Table 67. The correlation coefficients (R2) and power term (n) for the parabolic and log-log plots shown in Fig. 154 for isothermal oxidation of uncoated Ti- 22Al-27Nb. Temp. Parabolic log-log log-log t°C] (R2) (R } (n) 900 0.980 0.976 0.388 800 0.996 0.994 0.458 700 0.975 0.965 0.309 600 0.956 0.972 0.247 Table 68. The correlation coefficients (R2) and power term (n) for the parabolic and log-log plots shown in Fig. 155 for isothermal oxidation of the B-doped silicide coating grown on Ti-22Al-27Nb at 1150°c for 12 hours by a pack composed of 7 wt.% Si, 6% TiB2, 2% MgF2 and A1202. Temp. Parabolic log-log log-log [°C] (R2) (R2) (n) 1000 0.913 0.937 0.339 1000 0.906 0.846 0.565 900 0.963 0.876 0.043 900 0.979 0.967 0.106 800 0.911 0.956 0.251 800 0.901 0.856 0.421 700 0.922 0.945 0.378 600 0,978 0.974 0.534 600 0.996 0.989 0.744 500 0.895 0.885 0.334 ■ £ & (a) 1 2 3 4 5 6 Square Root Time [hour]1® «*r- 0.1^ B> E, — A—-500C *i 0-01 i - a- 600C CD - B- 600C £Z . .<>' 700C O) — X* 800C |0.001 n - . 0 . 900C — • © • 900C — 7 1000C D u —? 1000C 0.0001 -I— I ”1 I I 1111 T I 'I I I I I j 1- 1— I I I T ! I 0.1 1 10 100 Fioure 155. Linear regression fit for the isothermal oxidation data of a B-doped silicide coating Zr,v 22Al-27Nb by a pacX comprised of 7 wt.% Si, 6% TiB2, 2% MgF2 and A1203 (a) parabolic plot and (b) log-log plot. 350 -9— 700C -v - 600C -0 -* 5OOC 0.7 t — o - - 1000C 0.6 -x- • - 900C a - - 800C o a - *Q 00C o -©— 700C (3 0.3 Square Root Time (hour ]1 «g- l> E, c ’(9 0 sz O) 0.01 i 1 0.001 100 (b) Time [hourj Figure 156. Linear regression fit to the isothermal oxidation data for a Ge-doped silicide coating grown on CP- titanium at 1150°C for 12 hours by a pack composed of 12 wt.% Si, 6% Ge, 2% MgF2 and SiC (a) parabolic plot and (b) log-log plot. 351 0.25 ST 0 .2-3 £ 1000C *§> 900C 1,0.15^ A fi 2 o .h ** 600C 700C 800C ■§> '§ M s a a m m s ^ ™ > 0.05 TVjQPpT (a) Square Root Time [hour] 900C 1000C 0.1 i 600C E cd 0.01 - 0 700C ,g> | 0.001 0.0001 (b) Time [hour] Figure 157. Linear regression fit to the isothermal oxidation data for a Ge-doped silicide coating grown on CP- titanium at 1150°c for 12 hours by a pack comprised of 16 wt.% Si, 8% Ge, 2% A1F* and A1203 (a) parabolic plot and (b) log-log plot. 352 Table 69. The correlation coefficients (R2) and power term (n) for the parabolic and log-log plots shown in Figs. 156 and 157 for isothermal oxidation of the Ge-doped silicide coatings grown on CP-titanium at 1150°C for 12 hours by a pack comprising either (a) 12 wt.% Si, 6% Ge, 2% MgF2 and Sic or (b) 16% Si, 8% Ge, 2% A1F3 and A1203. Temp. Parabolic log-log log-log C°C] (R2) {R ) (n) (a) MgF2-Activated 1000 0.974 0.968 0.603 900 0.942 0.928 0.408 800 0.967 0.895 0.303 800 0.957 0.927 0.215 700 0.930 0.928 0.229 700 0.933 0.889 0.160 600 0.964 0.970 0.444 500 0.956 0.946 0.400 (b) A1F3-Activated 1000 0.967 0.947 0.106 900 0.891 0.929 0.072 800 0.954 0.982 0.202 700 0.972 0.958 0.223 600 0.897 0.795 0.096 500 0.996 0.983 0.873 353 1000C 0.2 5 - 0.2 i 0.15- 0 .0 5 - X ■800C 700C i i <*f | i rt i | i i i i | i i i ) | r-i i i | i i-i i | t i i i 1 2 3 4 5 6 : (a) Square Root Time [hour] 1/2 1000C 900C IT ■a o.i E, C ‘c5 0.001 I I 1 4 I 1 I 11 I 11 I I II I I 0.1 10 100 (b) Time [hour] Figure 158. Linear regression fit to the isothermal oxidation data for a Ge-doped silicide coating grown on Ti- 22Al-27Nb at 1150°C for 12 hours by a pack comprised of 12 wt.% Si, 6% Ge, 2% HgF2 and Sic (a) parabolic plot and (b) log-log plot. 354 Table 70. The correlation coefficients (R2) and power term (n) for the parabolic and log-log plots shown in Fig. 158 for isothermal oxidation of the Ge-doped silicide coating grown on Ti-22Al-27Nb at 1150°C for 12 hours by a pack composed of 12 wt.% Si, 6% Ge, 2% MgF2 and SiC. Temp. Parabolic log-log log-log t°C] (R2) (R ) (n) 1000 0.892 0.935 0.254 900 0.939 0.926 0.115 800 0.909 0.834 0.149 700 0.850 0.878 0.388 600 0.986 0.989 0.270 500 0.992 0.942 0.254 APPENDIX D X-RAY DIFFRACTION FOR BORON AND GERMANIUM-DOPED SILICIDE COATINGS FOLLOWING ISOTHERMAL OXIDATION F»« ■too- *0 - 70to " «0 * M * *0 * 90 * t10 o t i t a n i u m tnrctDt TTti* r i T m t w o h m / m i n x btm i.yro tyii > ,fn ■' 1,141* — 1 I.I 100 M| -• 00to • 00 TO * to • 40 .ta—A— /tA~ • *o BO >0 00 ■i 70j i i™r i 00 * o t i t am h q i o i u c i P t T J « t t TITAMIUM Q K IO t / IV T IL C . «TM t l - l t T I T J O t TITAMIUM OOOlOt Cb) Figure 159. B-doped silicide coating grown on CP-titanium at 1150°C for 12 hours by a pack comprised of 7 wt.% Si, 6% TiB-j, 2% MgF2 and A1203 following isothermal oxidation for 48 hours in air at (a) 900°C and (b) 600°C. 355 356 m 4 . 4 1 9 9. i . i. 1.199* 919.0*; EL £ L EL E l :EL -100 497.1 90 419.9* 90 191.3- 70 111.4' * 90 939.9* * 90 907,9 * 40 199.7* 30 103.9* 90 91.9 h io 0.0 i i I i i—i—i—j 11 i—n —i—i—i—i—i Tn i* i" i—i—i i i i 'i. i iA i* 0 90 30 40 SO 90 70 90 TITANIUM OXtOt t RUTtLE. 9TN 91-1979 ■ 1 Tt 09 TITANIUM SILICIDE 19-0799 -L. - U 1 i _ TI 919 ALUMINUM OXIDE / COOUWUM. SVN 10-0171 1 1 ..... jJ. (a) AL3 0 3 4.439 l.fTO E.pt l.fll 1.^41 l.}43 l.f99* 347.4- * 90 309.9 * 90 970.9* 70 911.9 90 193.0 * 90 1 9 4 , 4 - * 40 * 30 - 90 10 90 30 TITANIUM OXtOE / RUTILE. SYN 91-1179 TI 09 TITANIUM 9ILICI0E 19-0799 1 TI 919 TITAMIUM oonioe 39-0741 (b) TI 99 Figure 160. B-doped silicide coating formed on CP-titanium at 1150°C for 12 hours by a pack composed of 7 wt.% Si, 6% TaB-j, 2% MgF2 and A1203 following isothermal oxidation for 48 hours in air at (a) 800°C and (b) 600°C. :p b 4.439 9.979 9.£99 1 .f 93 l.^Al 1.743 1, 199,4- * 90 999.9* - 90 999.9* - 70 919.9* * SO 199.0* * 90 149.4* - 40 100.9* - 10 71.9* * 90 18.9* - 10 0 .0J I'll 1 ■|-|— n — i T'i— r~r- i— r 1 1 1 1 1 11 l"l I1 1 J I 1 1 1 1r 0 90 10 40 SO 90 70 90 TITANIUM oxloe / n u m e , stn 91-1979 1 . , | , 1. TI 09 ALUMINUM oxtoe / corundum. itn 10-0171 | I .. ll A AL9 01 NIOBIUM oxtoe 17-1499 ,, i J Jit It 1 It (A) N99 OS 4.m 191,0 J M . 7 - »0.4 994.1- ■ 17.* lit.9 149.9* I0I.9 71.1 ■0 90 40 TITANIUM oxtoe / BUTTLE. SYN ■1-1179 - 1 ^ - TI 09 ALIXIHUM OKIDg / COHUNOUN. SYN 10-0179 AL9 09 niobium oxtoe 17-1499 J_L J-*.. Jll.ll. J u 1 ■ '-i ■ 1 (b) NB9 09 Figure 161. Uncoated Ti-22Al-27Nb following isothermal oxidation for 48 hours in air at (a) 900°C and (b) 600°C 358 IP 8 4.416 2.f76 2.fS2 l.f23 1.J41 l.}43 1. 0 6 * 3 3 1 . 0 3 0 2 .4 - * 00 2 8 1 .8 * * 80 8 3 9 .2 - * 70 8 0 1 .6 - * 80 18 8 .0 * - SO 13 4 .4 * * 40 1 0 0 .8 - * 30 6 7 .2 * 1 fl > 1 6 - 20 3 3 . 6 - * 10 O .0 J 1"""« 1 1 "1 1 1 1 1 I 1 1 I'I'I I I 1 11 1 *V" 11 1 J 1 I'l 1 i-|" i"" - 0 SO 30 40 10 10 70 80 TITANIUM 1ILICI0E 10-0239 j __ L TI SIS TITANIUM 0*108 / MJT1LE. 1TN 11-1278 TI 02 ALUMINUM OXtOt / COaUNOUM. ITN 10-0673 (a) AL1 03 a»* <.«ie i.yia t.fsa i.fii t.yii t.yo i.iii* 867.1 sss.t 221.3 111.4 191,9 117.6 61.1 31.9 •0 >0 t i t a n i u m i n t e r n e 10-0213 J L. .1 i TI SIS TITANIUM OKIPg / RUTILE. IVN ti'ttrs TI 03 ALUMINUM OK IOC / C0RUN0UM. SVN 10*0173 A (b) AL1 03 Figure 162. B-doped silicide coating grown on Ti-22Al-27Nb at 1150°C for 12 hours by a pack composed of 7 wt.% Si, 6% TiB,, 2% MgF2 and A1203 following isothermal oxidation for 48 hours in air at (a) 900°C and (b) 600°C. 359 IPO 4 . 4 M 2 . f 7 0 ■ 2 3 2 l . f 2 3 l . f t l l . f 4 3 1 . 3 2 0 . 4 - - 0 0 2 0 2 . 0 - - 0 0 2 3 8 . 2 - - 7 0 2 1 0 . 0 - - 0 0 1 0 3 . 0 - • 3 0 1 4 8 * 4 - - 4 0 1 0 0 . 0 - - 3 0 7 3 . 2 - . 0 0 3 0 . 6 - ■ 10 2 0 3 0 4 0 ' ' i o ' ' 0 0 7 0 1 0 t i t a n i u m s i l i c i d e 3 3 - 0 7 8 3 1 J l l 1 ...... T I 6 1 3 TITANIUM OXIDE / M U TlLE. SVN 8 1 - 1 2 7 0 1 . 1 . 1 1 1. T I 0 2 SILICON OXIDE / CniOTOBALITE, SVN 3 0 - 1 4 2 3 ■I 0 2 1M. -100 * *0 - 00 310.O' - 7 0 ■ 7 3 . 0 ' - 00 ■ 1 7 , 0 ' * 00 ioi.O' 1 3 0 . 0 ' 30 01.0' ■0 40.0' 9 . 0 ' ■ 0 3 0 00 TITANIUM SILICIDE T I l i t TITANIUM OXIDE / WUTILC. QVN 1 1 - 1 3 7 0 T I 0 1 otLicoN oxiae t cbhtobalite. oym (to) ■ I 0 1 Figure 163. Ge-doped silicide coating grown on CP-titanium at 1150°c for 12 hours by a pack comprised of 12 wt.% Si, 6% Ge, 2% MgF;> and Sic following isothermal oxidation for 48 hours in air at (a) 900°C and (b) 600°C. a .a i b a.y7« a . » a i.pat i.y«t t. M t 100 B41.B an. * iBi,a 191.0 lao.o 90.B- BO.4-1 30.8* TITANIUM SILICIDE 39-07BS jL_ l J U u -L. Tt s i a TITANIUM OXIDE / ftUTlLC. BVN at-iars ti oa aluminum oxide / corundum, bvn 10*0173 alb oa 4 . 4 1 8 a . f7B a.yaa i.fti i.?4i 1 .^ 4 3 i . i100 m i 331.1 00 siB.a BO B71.1- * 70 B3B.4 BO 100.3- 30 130. B * 40 11B.T * 30 * to 30.07B.B 10 1 1 i 1 r r r r r T - r r BO 70 BO TITANIUM SILICIDE 19-07BS TI Bit TITAMIUM oxtoe / HUTILE. BYM I1-1B7B ti oa ALUMINUM OXIDE / COWUHOUM. 1VH 10-0173 (b) ALB 03 Figure 164. Ge-doped silicide coating formed on CP- titanium at 1150°C for 12 hours by a pack composed of 16 wt.% si, 8% Ge, 2% A1F3 and A1203 following isothermal oxidation for 48 hours in air at (a) 900°C and (b) 600°C. 361 4.436 a .pa a.pa i.yta i.£4 i i.pa *»i -•100 00 - 00 * 70 - *0 - so * 40 • 30 - 80 • 10 80 30 BO 00 00 TITANIUM 8IHCI0E t i cia TITANIUM 0KIQ8 / BUTILE. SVN TI 08 ALUMINUM CXI08 t CONUNOUH. BVN (a) ALX 03 3*0 4. £30 ■ ■£70 S. 1,£33 l.£41 1.J43 1,l»0* t a i i i a n • • « « t >•« n n - 00 - 00 ■ 70 • BO O ■ • BO • r-f j-r - 40 0 - 30 - 80 • 10 o.o- 1 1 II" SO 30 40 so 00 70 ■0 TITANIUM SILICIDC 10-0119 1 I l 1 1 i It ,11 TI (13 TITANIUM oxtoe / WJTILfc SYH i i ' t i n 1 . t , I t 1. TI 03 ALUMINUM 0X106 / COMMXJN, CVN 10-0173 1 1 _ AL3 03 Figure 165. Ge-doped silicide coating grown on Ti-22A1- 27Nb at 1150°C for 12 hours after 48 hours of isothermal oxidation in air at (a) 900°C and (b) 600°C. APPENDIX E CYCLIC OXIDATION DATA ■ Si-TiB2/%F2 r = — i- - Si-TaB2/AlF3 -•♦*-Si-TiB2/CuF2 —H - Si-TaB2/CuF2 — • • Uncoated CP Ti /<-Uncoated CP Ti 'E *§> £ oou V ae -c U «_* .2? jo (6.75 um) •f (2.6 um) <1'7 um) | iti 11 n 11II11m 11 ii 11111rri]111 40 80 120 160 200 (a) Number of Cycles Figure 166. Cyclic oxidation of B-doped silicide coatings grown on CP-titanium by six different packs at 1150°C for 12 hours with the thickness of the TiB? layer given in brackets. Coupons were tested at (a) 1000°C, (b) 900#C, (c) 7 0 0 ° c , (d) 6 0 0 ° c , and (e) 5 0 0 ® c . 362 Figure 166 continued. 1 2 Uncoated CP Ti (6.75 um) (3.9 um) (6.2 um) &!/■ (2.6 u m } (1.7 um) 0 ' ^ T T J ' i i i | i i ■ j ■ i i j"i I I | I I I | I I I | I I I | I j I | J I 1 0 40 80 120 160 200 i & - T . ’• T. . i .+•••+••+ ■ •- • • • I r (1.7 um 0* > I 1 I 1 1 1 I 1 I 1 I 1 1 1 1 1 I 1 40 80 120 160 200 (c) Number of Cycles 364 Figure 166 continued. 1.4rr (3.9 um) (6.2 um (1.7 um) (2.6 um) { <+. a r Uncoated C P' 'i I 1 l ■ l 1 l 1 I 1 I 1 I H ' H 1 40 80 120 160 200 (d) Number of Cycles 3 -TiB2/MgF2 *TaB2/^ftjF2 -TiB2/AlF3 -ToB2/AlF3“ 12/A J' 7 -TiB2/CuF2 _.-TaB2/CuF2 Uncoated CP Ti Q 9 um)^ "e ■8> £ Uncoated CP Ti u (2.8 um) bO cn j= (6 .2 um) \ (2.6 um) (6.75 um\^ 1,7 U\m^ U jz o n I A )n— O—'I 4> 1 * I 1 r ■ 1 1 r 40 80 120 160 (e) Number of Cycles 365 — — Si-TiB2/MgF2 1150C, 12h - •- - Si-TaB2/MgF2 U50C, 12h • • ♦ • • Si-TiB2/MgF2 950C, 6h - a - -SI/MgF2 U50C, 12h -•* -Uncoated CP Ti < • Uncoated CPU g Undoped Silicide (90*80 •S. g W B-doped (TaB O04 B-doped (TiB 8 0 1 2 0 (a) N u m b e r o f C y c le s Undoped Silicide (90-80 um) CM B Uncoated CP TI at E B-doped (TIB , 86-70 um) k _ i Figure 167. Cyclic oxidation of undoped and B-doped silicide coating grown on CP-titanium by a MgF2-activated pack at 1150°C for 12 hours or 950°C for 6 hours with the coating thickness given in brackets. Tested at (a) 1000°C, (b) 900°C, (c) 600°C, and (d) 500°C. Figure 167 continued. — •— Si-TiB2/MgF2 1I50C. 12h - •- * Si-TaB2/MgF2 1150C. 12h * • ♦ * • Si-TiB2/M gF2 950C, 6h - Si/M gF2 1150C, I2h - -x * * Uncoated CP Ti n ■ Undoped Silicide (90*80 um) ”s \i 6, B-doped (TiBj, 86-70 um) o 60 B-doped CT«B1P 75-65 «: = = H B-doped (TiB1, 25-20 um) . - -x- • -x* - *■"*' Uncoated CP *11 in - r | i i t | i i i ) i i i | i i i ) i i i | i i rj-ri i | i i i | i t i 0 40 80 120 160 200 (c) Number of Cycles 0 . 7 -q §lndoped Silicide (90-80 um) . ^ CM E f c w . u O) V - E r flidoped (TaB , 75-65 um) i _ i 0) D) L - - C (U '""T J—' - J= *- „ „ U B-doped (TlBz, 25|-20 um) +■» SI U) Tj-aopea 'a) 5 x - • x -x- -x - x- h v- • UncoatedCP l 1 I 1 I 1 I 1 i 1 i * i * i ’ i ’ 40 80 120 160 200 (d) Number of Cycles 367 Si-TiB2/AlF3 U50C, 12h - ■- • Si*TaB2/AlF3 1150C, 12h 50 — o- * Si/A1F3 1150C, 12h Uncoated CP Ti / f S 9 •a, Uncoated CP TI u , O' u Undoped Silicide (80-70 u m ) ^ | / DO B-doped (TaBj, 70-60 •s, -U — #3opedCnB}, 80-70 um i i11 I i iij i i i | i i i | i i i | i i i | i i i | i i i | r r r 80 120 160 2! (a) Number of Cycles N B-doped (TiB2, 80-70 um) E ■ U ncoated CP Ti O) E B-doped (TaB2, 70-60 i — i a> O) c ra x: CJ O) *a» 5 Undoped Silicide (80-70 um 1 11 * i i | I i i | i *1i ■ l^i I 1 ri-iyIM "W rrTi i i ^|~i lT i Wr | ■i ■i Ii | i i i 0 40 80 120 160 200 Number of Cycles (*>) Figure 168. Cyclic oxidation of undoped and B-doped silicide coating grown on CP-titanium by an AIF3-activated pack at 1150°C for 12 hours with the coating thickness given in brackets. Tested at (a) 1000°C, (b) 900°c, (c) 600QC, and Weight Change [mg/cm2] Weight Change [mg/cm2] (d) (c) .4 0 0.6 0.84 1.2 1.4J 1.6 0.2 .4 0 0.6 0.8 1.2 .4 1 1 0 1 JF I notdCU Undoped (80-70Silicide um) Uncoated CPU TB, 75-60 um) \ Uncoat?dCpT| T p C d ? t a o c n U \ ) m u 0 6 - 5 7 (TiBj,, d e p o d - B 0 6 1 0 2 1 0 8 0 4 I I I I 1 I ' I ’ I ’ I ' T ubr f Cycles of Number -Si Al 12h 2 1 , C 0 5 1 1 12h 3 , lF C 0 /A 5 1 2 1 B a T 3 i- lF S - /A 2 - iB * i-T S - •— — *Uncoat CP i T P C d te a o c n U * • * • - Cycl s le c y C f o r e b m u N U n d o p e d S ilic id e ( 8 0 - 7 0 u m ) ) m u 0 7 - 0 8 ( e id ilic S d e p o d n U doped (TaB, 70-60 um) ) m u 0 6 - 0 7 ,, B a T ( d e p o id B 0 2 1 0 8 / F3 1150C, h 2 1 , C 0 5 1 1 3 lF i/A S * S_ 2 'S _ B-doped um)80-70 (TiBj, B-doped (TaB2,*70-60um) -- ■ ------200 368 369 — — Si-TiB2/CuF2 1150C, I2h - *- * Si-TaB2/CuF2 U50C, 12h 35 t — a - *Si/CuF2 1150C, 12h “ • * * - Uncoated CP Ti n E 'AJncoatcd CP Ti / f I l Undoped Silicide (70*50 um)^ y / ! 00o B*doped (TaBj, 70*60 um) , ' X ^ ■a £ oped (TlBj, 70-60 um) ■ 1 ' i ' ' i1| * 11 | i 11 "i 111 j i 11 j 11 i | i 11 |" i i i | i i i (a) 0 40 80 120 160 200 Number of Cycles 2 0 B-doped (TiB , 11-8 um) N E Uncoated CPTi O) E Undoped Silicide (70-50 a> co c * * (0 JC i B-doped (TaB , 70-J50# urfO CJ JZ4 -1 CJ) *55 t- B-doped (TIB., 70-6(1 um) rr)"i i i |i i i | i ii | i i V| i 11 | 10T i i i i > 40 80 120 160 200 Number of Cycles Figure 169. Cyclic oxidation of undoped and B-doped silicide coating grown on CP-titanium by a CuF2-activated pack at 1150°C for 12 hours and 950°C for 6 hours with the coating thickness given in brackets. Tested at (a) 1000°C, (b) 900°C, (c) 600°C, and (d) 500°C. Figure 169 continued. 169 Figure Weight Change [mg/cm2] Wei8^ ° i2n^ f a g /c m 2] (d) 0.2 .4 0 0.6 0.8 1.2 1.4 0 1 If , 1, V • s i _ — ■* — ~ _ •si 2 10 0 0 2 160 120 0 8 0 4 0 rn'i'niTn i irn'i'niTn i i i i i i T i p r i i i n m ^■ ^ . i ■ r~r~ i ■ ~i » ■ i ■ i / / 0 0 2 0 6 1 0 2 1 0 8 0 4 K - • - • - K - - i • * - X - X * “ " V n c w & e d ” ^n| B-doped (TiBr 11-6um) B-doped (TiBr s e l c y C f o r e b m u N h 6 , C 0 5 9 2 F u C / 2 B i T - i S »•*Uncoated CP Ti T P C h d 2 e 1 t a o c , n C *U 0 • S ■» 1 - 1 2 F u /C 2 B i T - f S — — h 2 1 , C 0 5 1 1 2 F u C / i S • - * — h 2 1 , C 0 5 1 1 2 F u C / 2 B a T - i *S - - Undo] Cycl s le c y C f o r e b m u N i i > ■ d.»5 l ■ -*—*■■ »» ndop^Silicide (70-50 um) ndop^Silicide ) m u 0 5 - a ^ e d i c i l i S d e p o d n U dope TB 80-70 um 0 7 - 0 8 p (TiB ed p o -d 3 ___ B -d o p e d (TIB , 1 1 - 8 im ) ) im 8 - 1 1 , (TIB d e p o -d B B-doped um {TaBj, 70-60 B-doped ui (TiBJt70-60 . i » i- ■ 11 i"i ___ 371 — ■— Si-TiB2/MgF2 Ti-22Al-27Nb - - • - * Si-TiB2/M gF2 Ti-20AI-22Nb ...... Si-TaB2/MgF2 Ti-20Al-22Nb — x* • - Uncoated Ti-22Al-27Nb — h • U n c o a te d T i- 2 0 A l- 2 2 N b Uncoated Tl*22Al*27Nb _ _ * * • 1 -- 1.6 t>o TiBj, (20*22) 0.8 - •S) 0 . 4 - Uncoated Ti*20Al-22Nb 4 0 8 0 1 2 0 1 6 0 200 N u m b e r o f C y c le s 1.8 Uncoated TI-20AI-22Nb....« - « * • N E u T a B z , ( 2 0 - 2 2 ) s . 1.4* O) E 1.2 - j Uncoated U n c o a te d Ti-22AI-27Nb a) D) c TiBz, (20-22) TO s: 0.8 u o.6 -J ?J TiB2, (22-27) O) "5 0*4] 5 0.2 -p -r r |-i i i | i i 11 i i i | i i i [ i i i | i i i | i i i 40 80 120 160 200 (b) Number of Cycles Figure 170. Cyclic oxidation of B-doped silicide coatings grown on Ti-22Al-27Nb and Ti-20Al-22Nb by MgF2-activated packs at 1150°C for 12 hours, with the boron compound used in the pack given on the plot. Tested at (a) 1000°C, (b) 900°C, (c) 700°C, (d) 600°C, and (e) 500°C. iue 7 continued. 170 Figure Weight Change [mg/cm2] Weight Change [mg/cm2] of es le c y C f o r e b m u N ) d ( .5 0 0.6 .7 0 .4 0 .3 0 0 0 2 0 6 1 0 2 1 0 8 0 4 0 ••Si MgF2 -20Al b N 2 2 l- A 0 2 i- T 2 F g /M 2 B a T i- S • * *— - • * -Ti MgF2 -20Al b N b 2 N 2 7 l- 2 A l- 0 A 2 2 i- 2 T i- 2 T F 2 g F g /M 2 /M 2 iB T iB i- T S * i- * S • - - •— — — • ed -22Al b N 7 2 l- A 2 2 i- T d te a o c n U ■ • * • - Uncoated Ti-20Al-‘ Uncoated TiB, (22-27) TiB, h ed -20Al b N 2 2 l- A 0 2 i- T d te a o c n U * TiB, (20-22) TiB, TaB, (20-22) TaB, of es le c y C f o r e b m u N 120 Uncoated TI-22A1-27NIUncoated 160 200 iue 7 continued. 170 Figure Weight Change [mg/cm2] •; 4 . 0 .5 0 0.2 -E 3 . 0 0.1 ■: 0 Si MgF2 -20Al b N 2 2 l- A 0 2 i- b T N 2 b 2 F 2 g N - I 7 2 /M A 2 0 l- 2 B A a i- 2 T T 2 i- 2 i- S F T g 2 . F /M g . 2 . /M iB 2 . T i- . iB *S T . i- - . S • - - •— — ed -20Al b N 2 2 l- A 0 2 i- T d te a o c n U - * — 0 4 ed i Al b N 7 2 l* 2 A 2 i- T d te a o c n U of es le c y C f o r e b m u N 0 8 TiB, (22-27) TiB, 120 Uncoated Uncoated Ti-20Al-22Nb 0 6 1 200 374 — ■— Si-TiB2/MgF2 Ti-20Al-22Nb ---■*** Si-TiB2/CuF2 Ti-20Al-22Nb - ♦- * Si-TaB2/CuF2 Ti-20Al*22Nb - • * • • Uncoated Ti*22Al-27Nb — * - Uncoated Ti-20Al-22Nb 3 . 2 -i e t Uncoated Unco; Ti*22AI*27Nb ■&> cuFj, TiB, (20-22) E 2.4-j U>V C u F J( Ta B j (2 0 *2 2 ; nc 1.6-i j o " MgFj. 'HBj (20*22 .2? .. ■ > MgFj, *QBj (2 4 -11) . . j3 ; T , ' . * > ‘ V : '■ j .>• ' V- ~ ------— — m . Uncoated Ti*20AI*22Nb ¥■1«iT r i i | i i i | 111 p 11 | i i 11 i i i | 1111i 1 i |- m - 40 80 120 160 200 (a) Num ber of Cycles Uncoated Ti-20AI-22Nb CM E Uncoated Ti-22AI-27Nb o> E C uFz , T iB 0) O) c ra ■ I . . . I 11 i | il l I i i i | i i i | i i i | nr 1111 i j i i i 0 40 80 120 160 200 (*>) Number of Cycles Figure 171. Cyclic oxidation of B-doped silicide coatings grown on Ti-22Al-27Nb and Ti-20Al-22Nb by a CuF,- and MgF2- activated packs at 950°C for 6 hours, with the boron compound used in the pack given on the plot. Tested at (a) 1000°C, (b) 900°C, (c) 700°C, (d) 600°C, and (e) 500°C. iue 7 continued 171 Figure Weight Change [mg/cm2] i i " I1p l T I T l Uncoated Ti-20Al*22Nb Uncoated TI-22A1-27N Uncoated > Ti-20Al*22Nb Uncoated -TaB2/ Ti -22Nb N 2 2 l- A 0 2 i- T 2 F u b N /C 2 2 2 - B I a A T 0 i- 2 S i- - T 2 F g - /M 2 iB T - i- S •— — ed -20Al b N 2 2 l* A 0 2 i- b T N d 7 2 te a l- o A c 2 n 2 U i- - T d te a o x c n U * • — •x - 0 6 1 0 2 1 0 8 0 4 jiii|iiiji«»|iii jjiii|iiiji«»|iii of es le c y C f o r e b m u N -Ti CuF2 -20Al b N 2 2 l- A 0 2 i- T 2 F u /C 2 iB T i- S gFj. TiB gFj. C u F j . T a B j ( 2 0 - 2 2 ) F2 , T i B j ( 2 0 - 2 2 ) i T T j i 1 |’ i 375 Figure 171 continued. — •— Si-TiB2/MgF2 Ti-20Al-22Nb ...... Si-TiB2/CuF2 Ti-20AI«22Nb - - ■ Si-TaB2/CuF2 Ti-20Al-22Nb - * x • * Uncoated Ti-22Al-27Nb 0 .5 t “ X • Uncoated Ti-20Al-22Nb x-): Cu F2, TiB2 (20-22) V '6 Uncoated TI-20A1-22 £ MgFj, TiBj (20-22 o00 *s> ,^ X - x • X- UncoatedTI-22A1 »7Nb *------uFr TaBj (20-22) m ) 111 | i r r | i 11 [ ii n ' i u | > 11 |'< 11 |'n i |'H i 40 80 120 160 200 (d) Num ber of Cycles MgFr TiBj (22-27) , TiBj (20-22 E E & -X-Uncoated Ti-20A1-22N^>... «r;* S 0.05- | C ^ { £ .- - ■...... ’ ’ Uncoated Tl-22Al-27Nb ^ n n c 1 CuF„ TiB, (20-22) CuF,,TaB, (20-22 "U.Uj 11 i i i [t i i |*i i i I'vn | i n [ i 11 | i 11'| i M | iti | i rr 0 40 80 120 160 ( e ) N u m b e r o f C y c le s 377 ncoatcd Ti*22Al-27Nb (20-22: TaB2/CuF2-Activated) # / Ta •S s, 0 bO § #-i- - (20-22: ' -TiB,/CuF,-Activated — -G U i : ; - r » % f * 1 s' (20-22: TiB2/M gF2-AcUvated K U ncoated Ti-20Al-22Nb U ■ yiii|ii111111111[11i|11i|iii|ii11 i 11 | i i i 0 40 80 120 160 200 (a) Number of Cycles Figure 172. Cyclic oxidation of B-doped silicide coatings grown on Ti-22Al-27Nb and Ti-20Al-22Nb by a CuF2- and MgF2- activated packs at 950°C for 28 hours, with the boron compound used in the pack given on the plot. Tested at (a) 1000°C, (b) 800°C, (c) 500°C. Figure 172 continued. Uncoated Ti-20Al*22Nb Uncoated Ti-22Al-27Nb"* " 6 •Sb B TaBj/CuFj-Activated) § e (20-22: TiBj/CuFj-Activate*) § I (20-22: TiBj/MgFj-Activaied) 40 80 120 160 200 ( b ) N u m b e r o f C y c le s - ♦ •" .*(20-22: TaB , /CuF,-A ctivated) "e _ UncontedTj-20AU22Mb E x . Uncoated Ti-22A1-: ^ b . k- -x- - * § JG ' v (20-22: TiBj/MgFj-Activated) u JG«_> SP00 -0 .1 -i £ (20-22: TiB2/CuF2-Activateq) - 0.2 ■; " " ' ■ V I I I | I 11 | I I I | I I I | I I I | II If I I 17 I I I | I IT) I 40 80 120 160 200 Number of Cycles 379 — a— Si-Ge/A1F3/A1203 • • o • - Sl-Ge/MgF2/SiC — • • * * Uncoated CP Ti - •- • Si-Ge/CuF2/A1203 .Uncoated CP Ti Activa /CuF.-Acti rated u 15- ■ • a • O' C I Al ■ .0* OJ) s 10- r a** 8 ' « 6 * j - ° • ►^.0*'* • d^AlR-Activa cd •§)QU * 5 - I * CuFj-Activa^od — "* • I 1 1 1 I 1 I 1 I 1 I 1 I 1 I 1 40 80 120 160 200 Number of Cycles .UncoatedCPTi • _ A .. . . Je. . MgF-Activated t e 6 I 60u CuF2-Activated 80 120 160 Number of Cycles (b) Figure 173. Cyclic oxidation of Ge-doped silicide coatings grown on CP-titanium by an A1F3-, CuF,- and MgF,-activated pack at 1150°C for 12 hours. Tested at (a) 1000°C, (b) 900°C, (c) 700°C, (d) 600°C and (e) 500°C. Weight Change [mg/cm2] Weight Change [mg/cm2] continued. 173 Figure d Number of Cycles (d) c Number of Cycles (c) 0.8 0.2 0.4 0.6 0.5 4 8 10 6 200 160 120 80 40 0 UncoatpO CP Ti 40 Uncoated CP Ti 80 A1F-Activatj 120 -Activaled e l a v i t c A 2- F g M CuF-Activated CuF-Activated •Si-Ge/CuF2/A1203 Si-Ge/A1F3/A1203 160 yh CuF:-A itivated CuF:-A lFj-i1 ictivated >0* < ■ 200 380 iue 7 continued. 173 Figure Weight Change [mg/cm2] / !~' !~' •'"CuF,-Activated ^ / 0 0 2 10 200 160 120 80 40 2 A1F3-Acdvated 1 1 1 1 1 1 1 I 1 I 1 I 1t I 1 1 I I 1 I ° ' MgF-Activated ■°p ' Number Cyclesof • • -.•a-” CuF-Activated i** •• *o' - -Si-Ge/CuF2/AJ203 •- - i * •*• *— UncoatedCPTi ■• -• oSi-Ge/MgF2/SiC — Uncoated CP Ti b — — Si-Ge/A1F3/A1203 0 — — ..O *• 1 382 12Si*6Ge/Mg/MgF2/5iSiC o • * Uncoated CP Ti - - • 12Si*12Ge/MgF2/SiC - • * • 6Si-12Ge/Mgr2/SiC -X-- 10Si/MgF2/SiC 20-t 1:1 Ge*doped 5* ^ 0 15- ^ # / . j * - * 2:1 Ce*doped l:2Ge-dopedL \ ' rm ^ 1 \ . A ' a 10 Uncoatcd CP Ti •a £ jed Sillclde 0 | i-t i | i | i | i | i |-i | i l t (a) 0 40 80 120 160 200 Number of Cycles — 12SI-6Ge/MgF2/SiC Uncoated CP Ti - - • 12Si-12Ge/MgF2/SiC --•-•6Si-12Ge/MgF2/5iC —x--Si/MgF2/SiC Uncoated CP Ti 1:2 Ge-doped 1 1:1 Ge-do toU £ I Undoped; filicide , — ____ c-doped '<'1 ■ r r i ■ m i ' i 1 40 80 120 160 Number of Cycles Figure 174. Cyclic oxidation of Ge-doped silicide coatings grown on CP-titanium by a MgF2-activated pack at 1150°C for 12 hours with a Si to Ge ratio of either 2:1, 1:1 or 1:2. Tested at (a) 1000°C, (b) 900°C, (c) 700°C, (d) 600«C and (e) 500°C. 383 Figure 174 continued. 12Si-6Ge/MjjF2/SiCMgF2/, e • • Uncoated CP Ti - - - 12SI-12Ge/MgF2/SiCMs - - * - • 6Si-12Ge/MgF2/5iCAjf 1 —X -Si/MgF2/SiC f^.l:l Ge-doped 0,8 I 0'^~Uncoated CPTi 1 , •■1:2Ge-doped. a, 0.6 // * *• & I ° 7 '' «■' 1 0 . 4 t " 1:1 Ge-doped *f * ' ' '* * $ j ~ I °-2f .-V-]- ' -|-i | ■ i"| r |11 I | i | i | i [ I |~T 0 40 BO 120 160 (c) Number of Cycles — 12Si-6Ge/MgP2/SiC • • o • • Uncoated C PU - - - 12St-12Ge/MgF2/SiC - - • - • 6Si-12Ge/MgF2/SiCobi-UOe/Mgr - x - • • Si/MgF2/SiC 0 . 3 - i 1:2 Ge-doped „ 0,25 K"' p . - ■*’ Untfoped'Sil cide J:H 3e*d *--- ' Undoped S ific id e \ Uncoated CPTi i i ■ i » i 1 i 1 i 1 i ■ i 1 i 1 r 80 120 160 200 Number of Cycles iue 7 continued. 174 Figure Weight Change [mg/cm2] 4 . 0 t - 12Si MgF2/ iC /S 2 F g /M e G 2 1 i- S 2 1 • - - - 6Si MeF2/ iC /S 2 F e /M e G 2 1 i- S 6 • - • - - C i S / 2 Z F F g g M M / / i e S u 0 z 1 r - - - i - S x b • - — • - - • • * • o • • 1 Ge-doped e p o d - e G :1 2 0 6 1 0 2 1 0 8 0 4 ed te a o c n U -6Ge/ Si iC /S 2 F g /M e G 6 i- S 2 1 of es le c y C f o r e b m u N c i T P C / f b m ; d e p o d - e G Siici e id c ili S d e p o d n U ^ A 2 Gi d e p o d - i G ;2 1 384 385 16Si-8Ge/AIF3 Uncoated CP Ti 16Si-l6Ge/AlF3 8Si-16Ge/AlF3 x- ■ • Si/A1F3 Uncoated CP Ti G e - d o p e d o l:lJ3e-doped 6 P *&» E, u J Undoped SiHci w> JZ§ G e - d o p e d U +* xs an I Undoped Silicide I 1 I I '"I 1 I 1 I 8 0 1 2 0 Num ber of Cycles 16Si-8Ge/AlF3 • • o • • Uncoated CP Ti — . . . . 16SI*I6 G c/A1F3 - - • 8Si-16Ge/AIF3 1° ••••X— SI/AIF3 / Uncoated CPTi 4 8 4 / 1 : 1 Ge-dopei l U 6i ♦ /; j. * / < ^ 6 u 4-i • A - s * 1:2 Ge-boped $ I 2 x* Undgped Silicld^... i 'iy i ik i i * i i i i i i i i~ i ~\ i ~i i » fje-jinpr 0 40 80 120 160 200 f b ) Number of Cycles Figure 175. Cyclic oxidation of Ge-doped silicide coatings grown on CP-titanium by an A1F3-activated pack at 1150°C for 12 hours with a si to Ge ratio of either 2:1, 1:1 or 1:2. Tested at (a) 1000°C, (b) 900°C, (c) 700°C, (d) 600°C and (e) SOO^C. 386 Figure 175 continued. !6Si-8Ge/AlF3 • • o . . Uncoated CP Ti - - • 16Si-16Ge/AlF3 8Si-16Ge/AlF3 —X " I0SI/A1F3 : Uncoated CP Ti*, 1:1 Ge-doped-^.*- . —•*" _ 1 I B, 1:2 Gc-dopfflfl.i & ja t>0 gptf* 2:1 Cc-dopcd '£ ; ^ -K --U h'rfoped Silicide 0'T-i"'| i | 1 | 1 |"r | 1 | 1 | ■ -|- 1--|—i 40 80 120 160 200 (c) Number of Cycles — •— I6Si-8Ge/AlF3 . • o • • Uncoated CP Ti - - • 16Si*16Ge/AlF3 8Si-16Ge/AlF3 —X-SI/AIF3 e-doped 2 Gc- Ge-dogegL- Undoped Silicide 2:1 Ge-c I 80 120 Number of Cycles Figure 175 continued. 175 Figure Weight Change [mg/cm2] 1 Ge-doped e p o d - e G :1 2 - Al 3 lF /A e G i-8 S 6 1 — — * / F3 3 lF 3 lF lF i/A /A S e /A 0 G e 1 6 G •* 6 i-1 1 S i- 6 1 S x 8 • - — - • - - - - • • Uncoat CP i T P C d te a o c n U • • « • • of es le c y C f o r e b m u N 0 2 1 0 8 2 d e p o d - e G :2 1 d e p o d n U I lci e id ilic 387 388 — *— I6Si*8Gc/CuF2 . • o . • Uncoated CPU - - • 16Si*16Ge/CuF2 - • 8Si-16Ge/CuF2 —X 10Si/CuF2 llncoated CP Ti o 1:1 Ge*doped 1:2 Ge*dopeA' :1 Ge-dopcd 1 I 1 I 1 I T M I ' I 1 I 80 120 160 200 Number of Cycles — 16Si-8Gc/CuF2 • ■ o . . Uncoaled CPTi - - • 16Sl-16Ge/CuF2 - - • - • 8SU16Ge/CuF2 - x - • • Si/CuF2 Uncoated CP Ti 4 l:2Ge*doped u ^ \ ^Jpdoped Silicji / . ^ rfl Ge-dof «f (b j Number of Cycles Figure 176. Cyclic oxidation of Ge-doped silicide coatings grown on CP-titanium by a CuF2-activated pack at 1150 °C for 12 hours with a Si to Ge ratio of either 2:1, 1:1 or 1:2. Tested at (a) 1000°C, (b) 900«C, (c) 700°C, (d) 600°C and (e) 500°C, Figure 176 continued. 176 Figure Weight Change [mg/cm2] Weight Change [mg/cm7] 1*6- * Undoped Silicide '* 8Si-l6Ge/CuF2 ■• • • •Uncoated CP TI - - - 16Sl-!6Ge/CuF2 - - - 16Si-8Ge/CuF2 — — —x--Si/CuF2 Uncoated Uncoated CPTi 1:2 Ge-doped . 1:1Ge-dopcd 0 0 2 160 120 80 40 ' 1:5Ge-doppd •'* ^ 1:2Ge-doped, * 16Si*160e/CuF2 - - • - CP Uncoated Ti . •e • — ■— l6Si-8Ge/CuF2 l6Si-8Ge/CuF2 ■— — 8Si-16Ge/CuF2 x- • — • Si/CuF2 Number of Number Cycles of Number of Number Cycles of ----- ncoated * ' r f 120 160 200 200 160 120 JdJJejdoped - - * • 1 Un^spcd^lic^e ' 1:2 Ge-doped ■ 2:1Ge-do °* r0 ,mlr o * Ge-do )ed Ge-do 3 Figure 176 continued. 176 Figure Weight Change [mg/cm2] Number Cycl s le c y C f o r e b m u N ) e ( 1 Ge-doped e p o d - e G :1 2 ••*>•• • - - CuF2 2 2 F F u 2 u F /C u e /C G e i/C 6 S G 1 0 6 i- 1 l S - i 6 *• 1 S x* 8 • - — - 2 F • u - - /C - e G 8 i- S - 6 1 ■— — ed Ti T P C d te a o c n U : Ge-doped e p o d - e G 1:1 Sil cide d i c l i S d e p o d n U 391 — ■— Si-Ge/MgF2/SiC I150C, 12h U n c o a te d C P T i - *- • Si/MgF2/SiC 1150C, I2h — *••• Si*Oe/MgF2/SiC 950C,6h Uncoated CPTi E 1 u ISO •a (9 0 -8 0 u m ) JUntioped Silicide— (90-75* um) T IT i P ^I i | PI I Y*| I n I u I i i I 11 I i I r i I I I 11 40 80 120 160 2< (a) N u m b e r o f C y c le s Si-Ge/MgF2/SiC 1150C, 12h * • o * • U n c o a te d C P T i - »- - Si/MgF2/SiC 1150C, 12h — • ■ Si-Ge/MgF2/SiC 950C, 6h ~ 0 . 2 5 - Undoped Suicide (90*75 um) "<5(20-15 § 0 . 1 5 - (90-80 una) t e d C P T i Undoped Suicide (90-75 um) i • i • i 1 r 1 i"'1 i 1 i"'1 l"1 t 40 80 120 160 200 (b) Number of Cycles Figure 177. Cyclic oxidation of Undoped and Ge-doped silicide coatings grown on CP-titanium by a MgF2-activated pack at either 1150°C for 12 hours or 950°C for 6 hours with the coating thickness in brackets. Tested at (a) 900°Cf (b) 600°C, (C) 500WC. iue 7 continued. 177 Figure Weight Change [mg/cm2] fc) 05 .0 0 - 5 2 . 0 - 5 1 . 0 Sii de um) (90;80 ) m u 0 8 ; 0 9 ( ) m u 5 7 - 0 9 ( e id ilic S d e p o d n U -Ge/ Si 950C. . | h 6 . C 0 5 9 iC /S 2 F g /M e G i- S *Ge/ Si I2h h 2 I , C 0 5 1 1 iC /S 2 F g /M e G i* S — • — SiMgF2/ C 1150C, h 2 1 , C 0 5 1 1 iC /S 2 F g i/M S * - - * • ed Ti i T P C d te a o c n U • • o * • 0 0 2 10 : 160 120 80 40 Number of Cycles ""! » ! 1 I ■ 1 i IT"1" o. -o ed Ti T P C d te a o c n U 393 Si-Ge/A1F3 1150C, 12h a • • Uncoated CP Ti -* • • • Si*Ge/AlF3 950C, 6h 10 *■ -Si/A1F3 1150C, 12h U n c o a te d C P T i *¥ 8 f /(15-llum) JB \ « 6 & ?/' 4 Undoped Silicide (80-70 um > £ 2 40 80 120 160 N u m b e r o f C y cles — ■— Si-CJe/AlFi 1150C, 12h • ■ • • • Uncoated CP Ti - *- -Si/A1F3 1150C, 12h — *• • * Si-Ge/A1F3 950C, 6h .T-^-dS-ll ( 8 0 - 6 0 u m ) ■a Undoped Silicide (80-70 um ) I (8 0 - 6 $ u m ) d C P T i 40 80 120 160 Number of Cycles Figure 178. Cyclic oxidation of Undoped and Ge-doped silicide coatings grown on CP-titaniura by an AlF3-activated pack at either 1150°C for 12 hours or 950°c for 6 hours with the coating thickness in brackets. Tested at (a) 900°C, (b) 600°C, (c) 500°C. iue 7 continued. 178 Figure Weight Change [mg/cm2] (c) I I I I I I 1 I 1 I 1 1 I I 1 ' 1 I■ 1 I M 950C, ■ h 6 , C 0 5 9 3 F l A / e G - i S - h - 2 - 1 * , C 0 — 5 1 1 3 lF /A e G i* S — • — -SiA1F3 1150C, h 2 1 , C 0 5 1 1 3 F 1 i/A S - ♦ — • • o • . U n c o a te d C P T i i T P C d te a o c n U . • o • • 0 0 2 10 200 160 120 80 40 Number of Cycles l - 5 l f — 1 ) m u Silci (8p-70 ) m u 0 7 - p 8 ( e id lic i S d e p o d n U r'i't ed i T P C d te a o c n U 9 1' " T 394 395 • • o • - Uncoated CP H — ■— Si*Ge/CuF2 1150C, 12h — f ■ ■ - Si-Ge/CuF2 950C, 12h 10 - - * Si/CuF2 I150C, 12h Uncoated CP U 7— E 8 / (25*20 I y 6 / & / s // / S 4 / •' Undoped Silicide ^0*50 um) ^5) A / y S 2-3 (7 5 -6 5 u m ) 0 11111 n i 40 80 120 160 (a) N u m b e r o f C y c le s © • • Uncoated CP Ti ■*— Si*Ge/CuF2 1150C, 12h - * Si-Ge/CuF2 950C, 6h - - x - • Si/CuF2 1150C, 12b ( 2 5 -2 0 u m ) E ^ _ * - * - - - - - I u 60 / Undoped Silicide (70-50 um) J2OO (75-65 ur ) 0.2 4 r s Jndoped Silicide (70-50 um) a ,0.© _ . o ■ *o- • ® 0. . o. ° w Uncoated CPTi I 1 I 1 I ' I ' 1 ' I ' I ' 7 0 80 • 120 160 200 (b) N u m b e r o f C y c le s Figure 179. Cyclic oxidation of Undoped and Ge-doped silicide coatings grown on CP-titanium by a CuF2-activated pack at either 1150°C for 12 hours or 950#c for 6 hours with the coating thickness in brackets. Tested at (a) 900°C, (b) 600°C, (c) 500°C. 396 Figure 179 continued. — — Si-Ge/CuF2 1150C, 12h • ■ o • . Uncoated CP Ti — * *Si/CuF21150C, 12h — * ■••Si*Ge/CuF2 950C,6h "E *(25*20 u m f * 1 ■Si £ y r ' (75*65 um) & Undoped SiUcide (70 5 0 u m ) ( 2 5 - 1 0 u m ) .£P I .» • «• UnconijdCJTj,., 1 I 1 I 1 I I I I I ' I I I I I ' (o) 40 80 120 160 200 Number of Cycles 397 Uncoated Ti-20AI-22Nb CM CuF-Activated (2p-22) E o ! Uncoated Ti-22AI-27Nb z O) MgFz-Activated E AIFg-Activated (20-22) 0) O) c CO JZ U - r * o 4-> JZ ^ 4 ^ O) CD ,«%-*— • AIF-Activated (22-271) n 11 1111 r 11 | 1111 ii r | 11111 111111 111 111 40 80 120 160 200 (a) Number of Cycles Figure 180. Cyclic oxidation of Ge-doped silicide coatings grown on Ti-22Al-27Nb and Ti-20Al-22Nb by a MgF2-, AlF3-f and CuFj-activated pack at 1150°C for 12 hours. Tested at (a) 900°C, (b) 700°C, and (c) 600°c. Weight Change [mg/cm2] Weight Change [mg/cm2] 180 continued.Figure (b) (c) . U 0.1*1 2- .2 0 - .3 0 - .4 0 - .5 0 0.6* O-f v iV ! i i 111 r r | 11 | 11| ! i r i i i i 111 iV |i i 11^ i | r n r i v |i i11 O-f r ! 4 8 10 6 200 160 120 80 40 0 4 8 10 6 200 160 120 80 40 0 ► I * * < * * -" / I / i | i i i | i i i | S ^ ^ ^ •— — - * 5 2 . ) 2 2 - 0 2 ( " d e t a v i t c £ - ; F u £ i ed -Acti (£2-27) ) 7 2 - 2 £ ( d c t a iv t c A ,- F l A ) 7 2 - 2 2 ( d te a tiv c A - F g M ^ V -Activated ) 2 2 - 0 2 ( d e t a v i t c A ,- F u C Number of Cycles i -> Number of Cycles A c tiv a te d ( ^ 2 :2 7 ) , , ) 7 :2 2 ^ ( d te a tiv c A ii | i i i | rii | iM | i i i | i i r | i i i Uncoat Ti -22Nb N 2 2 l- A 0 2 i- T d te a o c n U ^ A 'x -Actvat (20-22) 2 2 - 0 2 ( d te a tiv c A ,- F 1 A ed - ' b N 2 -' 1 A 0 i-2 T d te a o c n U i ed ') 2 - 2 2 ( d te a tiv c A Ti22( - b N 7 ]-2 (A 2 i-2 T d e t a o c n U * J e27Nb N 7 2 te a o c n - 5 ! 398 399 — — Si*Ge/MgF2/SiC 1150C, 12h (22-27) - - * Si-Ge/MgF2/SiC 950C, 6h (20-22) - •» • - Uncoated Ti-22Al-27Nb — * * Uncoated Ti-20Al-22Nb Uncoated Tf-20AI*22Nb V | Uncoated T I-2 2 A I O) - - > E (20-22: 24-18 u a) a) c ra .c o O) (22-27: 150-120 nm) ’5 5 111 riT) i t i | i n"| i i'i'i'i i i | i i i | i i 11 iTT-p-i" 0 40 80 120 160 200 (a) Number of Cycles Uncoated Tl*20AI*22Nb ■&» E, Uncoated Ti*22Al*27 & (22*27:150-120t um) ^ (20*22:24-18 urn) •§» r . I ■ i j i • • j < i'i | i i r p - r r ( ii i | i i r p v ri1 '| ,it i 40 80 120 160 N u m b e r o f C y c l e s Figure 181. Cyclic oxidation of Ge-doped silicide coatings grown on Ti-22Al-27Nb and Ti-20Al-22Nb by a MgF2-activated pack at 950#C for 6 hours. Tested at (a) 900°C, (b) 700°C, -and (c) 600°c. Figure 181continued.Figure Weight Change [mg/cm2] (c) -12 H i - 4 - *8 0 » • ed -22Al b N 7 2 ) l- 7 A 2 - 2 2 2 2 i- ( T h 2 d 1 -* te , a o C c 0 5 n 1 U 1 • • ■» iC /S 2 F g /M e G i* S — * -Ge/ Si 950C,6h ) 2 2 - 0 2 ( h 6 , C 0 5 9 iC /S 2 F g /M e G i- S • - ed -20Al b N 2 2 l- A 0 2 i- T d te a o c n U - 0 8 0 4 ,(20-22:24-18 |im) ,(20-22:24-18 of es le c y C f o r e b m u N Uncoated 120 (22-27:150-120 Jim) (22-27:150-120 0 6 1 200 ■27Nb 401 *— Si-Ge/AIF3 1150C, 12h (22-27) - - • - • Si-6e/AIF3 1150C, 12h (20-22) - - - Si-Ge/AIF3 950C, 6h (20-22) -Uncoated Ti-22Ai-27Nb —* * Uncoated Ti-20AI-22Nb N Ut— Uncoated Ti-20AI-22Nb „ " E 0.5 i . I Uncoated Ti-22AI-2/Nb T O) \ ^ (20-22: 80-60 um) E, 0.4- a) i ' (20-22: 220 8 um) □> : I I c 0.3 i CTJ : I •C - v > - - O ~ 's' 4-> 0u *'-2- : ^ 4 SI .5* S I i • ■ i < ■ • 11111111 j 11111111111 |-i 111111 0 40 80 120 160 200 (a) Number of Cycles 0 .3 Uncoated Ti-20Al-22Nb - 6 JJnjoated Ti-22Al-27Nb B _____ ^6 . - r - - \ o r - <•-- ‘ (20-22: 81 eo c 0 . 1 5 - JSa U § — -vr £ 22:22-18 (un) 0-i tl 1 I | I I 1 | II I | ) I I j ! I I j ! '| I I I I I I ! I I I I I I | I I I 80 120 160 200 (b) N u m b e r o f C y c le s Figure 182. Cyclic oxidation of Ge-doped silicide coatings grown on Ti-22Al-27Nb and Ti-20Al-22Nb by an AlF3-activated pack at 950°c for 6 hours. Tested at (a) 900°C, (b) 700®C, and (c) 600°C. Figure 182Figurecontinued. Weight Change (mg/cm2) 0.2-j 0.5 0.6 0.1-i 0.4-j 0.3 -I A 4 8 10 6 200 160 120 80 40 0 i t -notd Ti-20AI-22Nb -Uncoated * — (20-22) 6h * 950C, Si-Ge/AIF3 - - (20-22) 12h - - -* • 1150C, Si-Ge/AIF3 - • * • - Uncoated Ti-22AI-27Nb Ti-22AI-27Nb •- • * - Uncoated (22-27) 12h 1150C, Si-Ge/AIF3 —— |ri| r / (20-22:22-18 um) (20-22:22-18 / 1 notdT-2l2N (22-27*. 100-80 um) Uncoated Ti-22Al-27Nb 1 111 | | 111 i Number of Cycles n | | n notdTi20At22Nb^^, ^ ^ b N 2 t-2 A 0 i-2 T Uncoated 11 i T | r ift-Tii | i p T|i i r i (20-22:80-60 jxm) (20-22:80-60 403 •Uncoated Ti-20AI-22Nb (VI I I (20-22: 100-80 oE N I* Uncoated Ti-22AI-27Nb O) li E i— j 0) O) c ro JC u *-» £ O) ’a) 5 (20-22: 13-9 jim) Number of Cycles Si-Ge/CuF2 1150C, 12h (20-22) - •- • Si-Gc/CuF2 950C, 6h (20-22) - • * • - Uncoated Ti-22Al-27Nb • Uncoated Ti-20Al-22Nb Uncoated Ti-2^1-22Nb ■¥ 100*80 um) j= ei Uncoated Ti •C op I I 11 111 l-j-l I I I I 11 I I I I I I I I I 11 11 I I I 40 80 120 160 200 Num ber of Cycles Figure 183. Cyclic oxidation of Ge-doped silicide coatings grown on Ti-22Al-27Nb and Ti-20Al-22Nb by a CuF2-activated pack at 9S0°C for 6 hours. Tested at (a) 900°C, (b) 700°c, and (c) 600°C. 404 Figure 183 continued. ■ Si*Ge/CuF2 1150C, 12h (20-22) - -Si-Ge/CuF2 950C,6h (20-22) • • Uncoated Ti-22Al-27Nb ~h * Uncoated Ti-20Al-22Nb 6 •S) Uncoated Ti-20AJ-22Nb^-- J3 y Si (20-22:100-80 um : 13-9 um) £ £ Uncoated TI-22A]-27Nb 80 120 Number of Cycles (C) REFERENCES 1. C.T. Sins, "Superalloys: Genesis and Character", in Superalloys II, C.T. Sims, N.S. Stoloff, and W.C. Hagel, (Eds.), John Wiley & Sons, New^York (1987) 3. 2. J.C. Williams, Proceedings of First International Conference on Structural Intermetallies, Seven Springs, PA, Sept. 26-29, 1993, R. Darolia, J.J. Lewandowski, C.T. Liu, P.L. Martin, D.B. Mitacle and M.V. Nathal (Eds.), TMS, Warrendale, PA, (1993) 839. 3. T. Van Aller, U.S. Patent # 1,155,974 (1911). 4. G.W Goward and L.W. Cannon, Trans. ASME, 110 (1988) 150. 5. M.G. Hocking, V. 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