A STUDY OF SELF-GLAZING

TITANIUM CARBIDE BASE CERMETS

Dissertation

Presented in Partial Fulfillment of the Requirements

for the Degree Doctor of Philosophy in the

Graduate School of The Ohio State

University

By

Robert Franklin Stoops, B.S., M.S.

The Ohio State University

1951

Approved by

Adviser ACKNOWLEDGMENTS

The author wishes to thank Mr. Earle T. Montgomery, Senior

Research Engineer, Mr. Thomas S. Shevlin, Research Engineer, and

Mr. Harold Greenhouse of the staff of the Ohio State University

Research Foundation Project 441 for their guidance, aid, and advice and for the design and construction of the equipment used in this investigation. He also wishes to express his appreciation to his adviser, Dr. G. A. Bole, Director of Ceramic Research,

Engineering Experiment Station, The Ohio State University, under whose general supervision this work was accomplished.

The work described herein was supported, in part, by the

Materials Laboratory, Research Division of Wright Air Development

Center through a contract between the Air Research and Development

Command and The Ohio State University Research Foundation.

ii

892567 The University assumes no responsibility for the accuracy or the correctness of any of the statements or opinions advanced in this dissertation.

iii TABLE OF CONTENTS

Page

INTRODUCTION ...... 1

SURVET OF LITERATURE...... 3

Highly-rRefractory Materials...... 3

Auxiliary Me t a l s ...... 7

Cermets...... 17

Sintering In the Presence of a Liquid Phase...... 20

MODE OF INVESTIGATION...... 23

MATERIALS AND EQUIPMENT...... 25

Materials...... 25

Equipment...... 30

PROCEDURE...... 39

Tests for Physical Properties...... 39

Titanium Carbide + Silicon Carbide Boron

Carbide Base Compositions...... 41

Titanium Carbide + Titanium Diboride + Silicon Base

Compositions...... 45

RESULTS AND DISCUSSION ...... 49

Titanium Carbide + Silicon Carbide + Boron

Carbide Base Compositions...... 49

Titanium Carbide + Titanium Diboride + Silicon Base

Compositions ...... 67

GENERAL S U M M A R Y ...... 97

iv TABLE OF CONTENTS

Page

CONCLUSIONS...... 99

BIBLIOGRAPHY ...... 101

AUTOBIOGRAPHY ...... 105

v ILLUSTRATIONS

Figures Page

1. Ratios of Specific Volumes of Oxides of Metals To

Those of the Metals...... 11

2. Phase Diagram of the Cobalt-Silicon System...... 16

3. Molybdenum-Resistor Furnace ...... 32

4. Section Through Center of Molybdenum-Resistor Furnace . . 33

5. High Frequency Induction Furnace...... 34

6. Cross Section of High Frequency Induction Furnace.... 35

7. Globar Tube Furnace...... 37

8. Hot Modulus of Rupture Furnace...... 38

9. 721 R + Co Sintered at 3875°F. (Photomicrograph)...... 53

10. 721 R Jr NiAl Sintered at 3375°F. (Photomicrograph). . . . 54

11. 721 R Sintered at 4100°F. (Photomicrograph)...... 58

12. 721 Sintered at 4100°F. (Photomicrograph) ...... 59

13. Effect of Sintering Temperature on Porosity of

TiC + SiC + B.C Base Cermets...... 63 4 14. Effect of Sintering Temperature on Porosity of

TiC + TiB2 Si Base Cermets...... 72

15. Composition III R Co Sintered at 2800°F.

(Photomicrograph) ...... 74

16. Long Time Oxidation in Open Air at 2000°F...... 77

17. Effect of Changes in Composition on Oxidation of

TiC + Ti®2 + Si + Co Cermets...... 83

vi ILLUSTRATIONS

Figures Page

IS. Effect of Sintering Temperature on Transverse

Strength of TiC + T±B^ 4- Si 4- Co Cermets...... 87

19. Graphite Boats Containing Fired B a r s ...... 89

20. Ill R B + Co Sintered at 2800°F. (Photomicrograph) .... 91

21. Change of Modulus of Rupture of Composition

III R B +• Co with Temperature...... 94

Tables

I. Data on Highly-Refractory Materials...... 26

II. Particle Sizes of Milled Carbides...... 28

III. Data on Raw Materials...... 29

IV. Compositions Studied...... 46

V. Results of Wetting Tests ...... 52

VI. Results of Sintering Tests on 721 Base Composition .... 56

VII. Porosities of Sintered Cermet Pellets Based on 721 .... 62

VIII. Effect of Sintering Temperature on Properties of

721 R 4- C o ...... 65

IX. Porosities of Sintered Cermet Pellets Based on

TiC 4- TiB2 + Si...... 69

X. Results of Oxidation Tests ...... 75

XI. Values for the Parameter "k" from Long Time

Oxidation Tests...... 76

vii ILLUSTRATIONS

Tables Page

XII. Effect of Sintering Temperature on Properties of

TiC ♦ TiB£ 4- Si + Co Cermets...... 86

XIII. Change of Modulus of Rupture of Composition

III R B + Co with temperature...... 93

XIV. Results of Thermal Shock Tests...... 95

viii INTRODUCTION

The efficiency of gas turbines would be greatly increased if

they could be operated at temperatures above those currently being

used. The metallic alloys from which the turbine blades are made

have a relatively short service life at an operating temperature of

1800°F. At present, a material is being sought which will have a

useful service life at a temperature between 1800°F. and 2400°F.

Such a material should have a high tensile strength to density ratio

in this temperature range because the stresses developed in turbine

rotor blades are caused mainly by centrifugal forces. The material

should also be resistant to oxidation and to thermal and mechanical

shock.

Much research has been directed toward combining highly-refrac-

tory ceramic materials with metals in order to obtain a material that

possesses both the high strength to density ratio at elevated temper­

atures of ceramics and the thermal and mechanical shock resistance

of metals. Some of these combinations, called "cermets", which show promise utilize titanium carbide as the ceramic phase. However, these cermets lack the necessary oxidation resistance, since their

chief constituent, titanium carbide, oxidizes readily at elevated temperatures. If these compositions are to be used successfully in turbine blades, they must be coated with an oxidation resistant material or other constituents must be added to the cermet composi­ tions which will render them oxidation resistant. In an investigation of the titanium carbide - silicon carbide - boron carbide system, Accountius^ found several compositions rich

in titanium carbide which had excellent oxidation resistance. These materials were oxidation resistant as a result of the formation of a glass on their surfaces when they were exposed to oxidizing condi­ tions. The purpose of this investigation is to develop a self-glaz­ ing cermet by using one of these oxidation resistant compositions as a base material. SURVEY OF LITERATURE

HIGHLY-REFRACTORY CERAMIC MATERIALS

Schwarzkopf^ has stated that materials will have high strength at elevated temperatures if the individual crystals have high strength interatomic bonds and if the various crystals have high boundary strength. Highly-refractory materials possess relatively strong inter­ atomic bonds since the energy requirements for melting are related to the atom-to-atom bond strength. He suggests that the powder metal­ lurgy technique could be utilized to control boundary strength.

Titanium Carbide

Titanium carbide is one of the. most refractory materials known.

It has a melting temperature of 5684°F. Gangler, Robards, and 17 McNutt ', who investigated the physical properties at elevated temper­ atures of hot-pressed TiC, MgO, ZrC, B^C, 85$ SiC 15$ B^C, ZrOg, and stabilized Zr02, found that titanium carbide had the best resistance to thermal shock and was generally the most promising of the composi­ tions tested. It had a short-time tensile strength of 15,850 p.s.i. at 1800°F. and 9400 p.s.i. at 2200°F. Deutsch, Repko, and Lidman^ concluded from their investigation that titanium carbide base cermets may eventually be used as gas turbine blade materials in the temper­ ature range of 1600°F. to 2400°F. 33 Nowotny and Glenk asserted that titanium carbide tends to have a lattice which is deficient in carbon. The best titanium carbide

-3- that they could obtain had only 19*36# carbon, whereas the theoret­

ical value is 20.0#. Franssen"^ has stated that, as a rule, com­

mercial titanium carbides do not contain more than 18# of fixed

carbon; but, in order to obtain a fixed carbon content of 18#, an

excess of carbon must be used. This free carbon remains in the

carbide as graphite. Skaupy^ also thinks that commercial titanium

carbides contain less than theoretical carbon, but that they contain

an oxide of titanium in solid solution in the titanium carbide. This

oxide of titanium comes from the process in which rutile and carbon

are combined to make titanium carbide. Such a carbide containing

oxygen will not produce dense sintered compacts. Skaupy believes

that the carbide made by Kennametal, Inc., is superior to other

commercial titanium carbides because it contains little or no 31 oxygen. Meerson and his associates have explained the decrease m

the density of impure titanium carbide as bloating caused by gas

evolution.

Silicon Carbide

Silicon carbide dissociates into silicon and carbon at 3992°F.^ without melting. Sabol-^ was unable to obtain dense compacts when he attempted to sinter combinations of alumina and silicon carbide, although dense bodies were obtained when alumina was sintered alone.

Weight losses during the sintering of the combinations were attributed to the formation of volatile silicon monoxide and possibly hydrides, although the firing temperature range was 3300° to 3500°F. Lea^ has said that silicon carbide is of interest not only because it is refractory, but also because it has good oxidation resistance, high hot strength, and good thermal shock resistance. Lefebvre^ found that silicon carbide oxidized slowly up to 3362°F., but that above this temperature the protective film of silica melted away.

Boron Carbide

Boron carbide melts at 4262°F. and has a specific gravity of 3 17 2.52 . Gangler, Robards, and McNutt 1 have reported that boron carbide has a short-time tensile strength of 22,550 p.s.i. at

1800°F. They said that this relatively high tensile strength and the low specific gravity of boron carbide give this material a high strength to weight ratio that is important for bodies that are subjected to centrifugal forces. The main disadvantage of boron carbide is its poor thermal, shock resistance. 28 Lidman and Hamjian have asserted that, during oxidation, boron carbide forms a glassy surface coating of boron oxide that sublimes in the temperature range from 1800° to 2400°F. Lefebvre^ said that this material oxidizes only slightly above 1832°F.

Oxidation Resistant Titanium Carbide Base Composition

In examining compositions containing titanium carbide, silicon carbide, and boron carbide, Accountius^ found several oxidation resistant compositions which contained ten percent boron carbide, between fifty and seventy percent titanium carbide, and between

-5- forty and twenty percent silicon carbide. When the sintered mix­ tures of the carbides were examined by x-ray and metallographic methods, they were found to consist of titanium carbide, titanium diboride, material of the composition designated as TiBx, silicon carbide, and graphite produced from the reaction between the titanium carbide and boron carbide. Metallographic examination revealed that these bodies were not homogeneous. Their cross sections consisted of two different types of areas. One type of area was rich in titanium carbide and contained a lesser quantity of titanium diboride, while the second type of area contained all of the silicon carbide, and considerable amounts of titanium carbide, titanium diboride, and graphite. The compound TiB„ was not located.

Accountiu3^ also reported that when the sintered bodies of the oxidation resistant compositions were oxidized in air at 20C0°F. and their oxide coatings were investigated by x-ray, metallographic, and petrographic methods, the coatings were found to consist of titanium dioxide, silica, and glass with an index of refraction of 1.52.

This was thought to be a boro-silicate glass which contained some titanium dioxide. The most oxidation resistant composition contained a weight ratio of silicon carbide to boron carbide of approximately two to one.

Titanium Diboride

The melting temperature of titanium diboride was not found in the literature, but Norton and his associates^ consider it to be

-6- a very refractory material with well developed metallic properties.

The electrical and thermal conductivities of this material compare

favorably with those of pure metals.

The oxidation of titanium diboride was studied at 2192°F. by

Tinklepaugh^*", and this material was found to be much more resistant to oxidation than titanium carbide, even when the fact was considered that there was probably some volatilization of the oxide layer. g Bennett and others have also reported a reaction between the tita­ nium carbide and boron carbide to form titanium diboride and carbon.

The reaction began as low as 3000°F. and was complete after sinter­ ing of the carbide mixture at 3750°F* for one-half hour.

Titanium Disilicide

Honigschmid reported that titanium disilicide crystallizes in iron gray tetragonal pyramids, and that it has a hardness of 4 or 5 and a specific gravity of 4.02 at 22°C. He also stated that this silicide has very good resistance to oxidation.

AUXILIARY METALS

Skaupy^ found that the mechanical properties of a cermet with a definite carbide as base material depend to a large extent on the nature of the auxiliary metal used. Even metals which are as nearly alike as cobalt and nickel may show wide differences in behavior when used as binding or auxiliary metals. Redmond^ states that the binder metal must be one in which the carbide is soluble in reasonably high percentage at the sintering temperature, yet which has little retention for the carbide at ordinary temperatures. If the metal retains the carbide in solution, the metal will be embrit­ tled and the amount of metal will be effectively increased, accord­ ing to Skaupy^". A metal which has the most desirable type of solubility for a carbide will dissolve the carbide during sintering and precipitate it during cooling. He thinks that no new carbide nuclei are formed, but that all of the precipitation occurs on grains already present. In this way, a continuous carbide skeleton is formed unless there is so much metal present that it more than fills the interstices between the carbide grains and holds them 36 apart. Redmond-3 has verified the existence of skeletons formed in this manner by leaching the auxiliary metal with acid. 12 Dawihl and Hinnueber report that binder metals which meet the above solubility requirements reduce the sintering temperature necessary to obtain optimum properties from a carbide. For example, the temperature at which the optimum stress-rupture strength is attained is reduced from about 3270°F. for tungsten carbide alone to about 2370°F. for tungsten carbide containing 11% cobalt. They also said that metals which do not dissolve the carbide phase will obstruct sintering and increase the sintering temperatures neces­ sary to obtain optimum properties. 36 According to Redmond-^ , auxiliary metals should be ductile to some extent, and they should have little or no affinity for carbon so that they will not decompose the carbide phases. Deutsch, Repko,

-8- 13 and Lidman found that the more refractory metals impart higher strengths in the temperature range of 1600° to 2400°F. than less refractory metals.

Oxidation Of Metals

Deutsch, Repko, and Lidman^ think that it is important to use a metal that is oxidation resistant since the oxidation characteris­ tics of cermets are altered by the auxiliary metal added. Pilling •ac and Bedworth-^ state that noble metals are oxidation resistant because at temperatures below their melting points they form oxides having oxygen pressures which are higher than the partial pressure of oxygen in the atmosphere. Base metals are oxidation resistant if they form an oxide coating which inhibits further oxidation. 1 15 Evans * has given the following three basic requirements that the oxide of a metal must meet before it will inhibit further oxidation: first, the volume occupied by the oxide formed must be greater than that originally occupied by the metal oxidized so that a dense film of oxide will be formed; second, the oxide must adhere tenaciously to the unoxidized metal, and, finally, the oxide must be of such a nature that neither the metallic ions nor oxygen ions can migrate through it.

According to Pilling and Bedworth^, if the volume of the oxide formed is less than that of the metal which it displaces, a porous coating is formed through which oxygen can easily pass to unite with the metal surface. If the oxide film formed on a metal is dense,

-9- further oxidation can take place only by diffusion of ions and/or

electrons through the crystal lattice of the oxide. They use the following ratio to determine whether or not the volume of a metallic oxide is greater than that of the metal it replaces:

M d m D where M is the molecular weight of the oxide

m is the molecular weight of the metal

D is the density of the oxide

d is the density of the metal.

If this ratio is less than unity, a porous, unprotactive oxide coat­ ing will be formed? but, if this ratio is equal to or greater than unity, a dense protective film may be formed. Figure 1 shows the values of this ratio for various metals as given by Lustman2^. The metals which have been found to form tightly adhering oxides are underlined.

Pilling and Bedworth-^ believed that, when dense oxide coatings were lormed, further oxidation could occur only by diffusion of oxygen inwards toward the metal surface. Such ionic movement would occur directly through the crystal lattices. Evans^ has reported more recent experiments which have shown that in many cases the metallic cations and the electrons diffuse outward through the dense oxide coating to meet the oxygen. According to this theory, oxida- tion can take place only if the oxide film is capable of conducting both metallic ions and electrons. An outward movement of the

-10- Li Be 0 3 7 1.59 Klq Mo Al S i 031 0.79 1.282.27 K C q T i V Cr Mr! Pe Co Hi 0 4 6 0.65 1.95 8.18 1.99 1.79 177 1.99 1.52 Cu Zh Ge A s (.68 1.62 1.64 2.15 Rb S r Zr Cb Mo 0.450.65 !.5| 2.61 3 .4 1.60 Aq Cd Ih Sr> Sb 1.5% I.ZI 128 1.32 2.85 Cs B q Lq Ce W 0.460.74 111 1.16 3.4

H p To • 1.48 2.33 h 9 Tl Pb Bi 1.81 1.88 1.40 2.27 Th U 1.32 3.05

Figure 1.

Ratios of specific volumes of oxides of metals to the specific volumes of the metals as given by Lustman29. The underlined metals have been found to form tenacious, protec­ tive oxides in hot, oxidizing atmospheres. These metals are arranged in periodic sequence.

-11- cations would be favored by vacant metallic atom positions in the

oxide. Evans stated that the composition of the oxide of a metal of variable valency often differs appreciably from that suggested by the formula. Since these "metal deficient" oxides are likely to be good conductors of electricity, they do not protect the metal from oxidation. Thus, the electrical conductivity of the oxide is sometimes used to predict the protective ability of the oxide. 29 Lustman states that, when non-protective oxides are formed, oxidation will take place at a rate that is usually constant. How­ ever, when a protective coating is formed, oxidation will normally proceed according to the following equation:

W2 = k t where W is the thickness or weight of oxide scale formed

k is a rate constant

t is the time.

After a sufficient length of time, oxidation which follows this relation should be negligible. However, as the oxide coating grows in thickness, the possibility that it will peel or crack becomes correspondingly greater. Therefore, Pilling and Bedworthr^ have suggested that it is very desirable to utilize a material which forms an oxide so impervious that even a very small thickness is an almost perfect barrier to oxidation.

Effect Of Percentage Of Metal On Skeleton Structure

13 Deutsch, Repko, and Lidman report that auxiliary metals are

-12- usually,incorporated in a cermet composition as fine powders in quantities that vary from several percent to twenty or thirty percent, depending on the use for which the material is intended. When the metal content is low, the metal acts as a catalyst for the formation of a strong ceramic skeleton. At higher metal contents, the metal tends to bond the ceramic particles into a continuous metallic net­ work. In an investigation of tungsten carbide cermets Skaupy^ found that when the cobalt content was approximately three percent or less the carbide skeleton remained incomplete and that the toughness of the resultant cemented carbide was insufficient for many uses. He also observed that in the second range, from three to about eight percent of cobalt, the carbide skeleton formation was complete as was shown by the increased toughness of the result­ ant material. As would be expected, even small amounts of cobalt had a beneficial effect on the cermets in that they filled the pores of the carbide and helped to absorb local stresses. When the cobalt content was more than eight percent, Skaupy found that the skeleton was again not complete because the excessive quantity of cobalt held the carbide grains apart.

Cobalt

Commercial cermets usually contain a metal of the iron group such as cobalt or nickel, according to Skaupy^1. Nowotny and Glenk33 state that cobalt has been found to be superior to iron and nickel for use in titanium carbide base cermets because it will not retain

-13- titanium carbide in solution at low temperatures, yet it will

form a solid solution with titanium carbide at elevated temper­

atures. The solubility of titanium carbide in cobalt has been

reported by Meerson and others^ to be from seven to ten percent

in the temperature range of 2100° to 2280°F. Cobalt is more

ductile than other common binder metals, and it possesses good

strength, according to Rose^®. Hamjian and Lidman^® have bonded

boron carbide with cobalt and Heyroth^ has bonded silicon

carbide with this metal.

. Nickel

Konrad and Stoops^ and Wulff^ have used nickel as a

binder metal with titanium carbide. Boron carbide was bonded

with nickel by Hamjian and Lidraan^, and HeyrothA^ succeeded in

wetting silicon carbide with nickel at a temperature of approx­

imately 3632°F.

Chromium

18 Hamjian and Lidman found that molten chromium metal wetted

boron carbide under suitable firing conditions. Titanium carbide was also wetted by chromium in experiments conducted by Konrad 21 and Stoops .

Nickel-Aluminum Intermetallic Compound

31 McBride used the intermetallic compound nickel-aluminum as a bonding phase for titanium carbide, but he did not obtain

dense bodies of this composition. Nevertheless, these cermets

had good strength at elevated temperatures.

Cobalt Silicides

The various authors disagree as to how many different cobalt-

silicon compounds exist, and also as to their physical properties.

Figure 2 is a phase diagram of the cobalt-silicon system as given p by Hansen*. This diagram conforms, in general, to the opinions

held by the majority of the authors. The melting points of the

compounds and the eutectic and peritectic temperatures within

this phase diagram have slightly different values in the various

references.

Baraduc-Muller^ stated that cobalt hemisilicide, Co2Si, forms

as steel gray crystals that are very hard and brittle. Boren,

Stahl, and Westgren^ have found that the crystal structure is

orthorhombic, and they have given the unit cell dimensions,

atom positions in the unit cell, and x-ray data. Boren^ has reported that cobalt hemisilicide is stable only below about

1832°F. This silieide has a specific gravity of 7.1 at 17°C., according to Mellor^.

Baraduc-Muller^ gives 6.30 as the specific gravity of cobalt monosilicide, CoSi, at 20°C. He said that this silieide had been reported as being not very hard nor very brittle and that it dissociated into cobalt hemisilicide and silicon at high temper-

-15- 1500 - .144 Cf

1327

1277

2 3 6 ' 3 7J^I l20l I h 120°--- - 3 =

2 1100

1000

900 Co lO zo 30 4 0 50

F ig u re 2 . 23 atures. Lebeau found that cobalt monosilicide had remarkable

oxidation resistance and that it formed brilliant crystals which

belonged to the cubic system, A solid state reaction between this

silieide and cobalt hemisilicide to form the compound Co^Sig has been reported by Baraduc-Muller^ and Lewkonja^.

Lebeau^- found that cobalt disilicide was made up of dark, blue-gray crystals, which were thought to be cubic and to present octahedral form most frequently. Baraduc-Muller reported that this

silieide was brittle and had a specific gravity of 5.3 at 0°G. He also stated that cobalt disilicide can be separated from the mono­

silicide by dissolving the latter in hydrochloric acid.

CERMETS

Titanium Carbide Base Cermets

31 Meerson and his coworkers^ found that wetting of the titanium carbide grains by cobalt was obstructed by the evolution of carbon monoxide when commercial titanium carbide of low purity was used.

They thought that this same gas evolution produced porosity in the sintered bodies.

Whitman and Repko^ reported that when titanium carbide was milled with either molybdenum, tungsten, or cobalt an appreciable amount of tungsten carbide was picked up by the materials from the tungsten carbide balls used in milling. When cobalt was milled with titanium carbide, usually less than one percent but sometimes as high as 2.1 percent of tungsten in the form of tungsten carbide was present as an impurity.

-17- In an investigation of a cermet composition containing eighty- percent titanium carbide and twenty- percent cobalt, Hoffman, Ault, and Gangler^ produced bodies which had a short-time tensile strength of 33,200 p.s.i. at 1800°F. and as high as 13,200 p.s.i. at 2200°F.

This same composition had modulus of rupture values which varied from 57,800 p.s.i. to 100,800 p.s.i. at 1600°F. and which ranged from 29,400 p.s.i. to 71,900 p.s.i. at 2000°F. Whitman and Repko^* stated that titanium carbide bonded with cobalt possesses many of the properties required for a material to be used in gas turbine blades, but that it lacks sufficient oxidation resistance.

In an investigation of mixtures of titanium carbide and boron C> carbide bonded with cobalt or nickel, Bennett and others0 found the compound cobalt monoboride in all of the sintered bodies which contained both cobalt and boron carbide, and they found nickel borides in all sintered compositions containing nickel and boron carbide. Neither of these borides showed any tendency to bond the titanium diboride. They also reported the reaction between titanium carbide and boron carbide to form titanium diboride and free graphite. None of these cermets containing boron carbide were as strong as titanium carbide alone bonded with the metals.

Whiff7*8 investigated a cermet containing 80$ titanium carbide and 20$ nickel. He found that the carbide produced by Kennametal,

Inc. produced denser bodies than did the titanium carbide made by

Titanium Alloys Manufacturing Division, National Lead Company. The two carbides were made by different processes and contained

-18- different amounts of impurities. Wulff milled his cermet composi­ tion for twenty hours in carbon tetrachloride. He sintered the compacts at 1100°C. for one hour in a vacuum before the final sin­ tering because denser specimens were obtained as compared to those produced without the pre-firing operation. He attributed the increased density to the elimination of gas pockets. The titanium carbide - nickel cermets were densest when sintered for one and a half hours at 1400°C. The modulus of rupture values varied from 80,000 p.s.i. to 130,000 p.s.i. with the most consistently high values resulting from sintering at 1350°G. for four hours.

Boron Carbide Base Cermets

18 Hamjian and Lidman conducted bonding experiments with boron carbide and cobalt, nickel, and iron. These tests indicated that, under proper firing conditions, a bonding phase forms between each of these metals and boron carbide. The liquid metals selectively dissolve part of the boron carbide and penetrate into the pores to form the bonding phases. A boron carbide - iron cermet body had good strength properties at 2600°F., and this fact led Lidman and 18 28 Hamjian * to postulate that materials that form a bonding phase may possess desirable strength properties at temperatures near the melting points of the auxiliary metals.

-19- SINTERING IN THE PRESENCE OF A LIQUID PHASE

Furnace Atmosphere

Cermet compositions are usually sintered in neutral or reducing atmospheres to prevent oxidation and decarburization, according to

Rose^®. Hydrogen is usually used as the atmosphere, but compacts containing titanium carbide are sometimes sintered in a vacuum. 37 Rhines stated that an atmosphere which will reduce surface oxide films will generally improve the strength and density of cermets.

Effect Of Time On Sintering

37 Rhines reported that the rate of sintering is most rapid at the beginning of the sintering treatment, and that the effect of time is most noticeable at low sintering temperatures. He suggests that long sintering times may cause more lowering of strength associated with grain growth than increase of strength resulting from shrinkage 38 of the body. Rose has said that cemented carbides are usually sintered from a half-hour to an hour.

Mechanism of Sintering In The Presence Of A Liquid Phase

According to Lenel^, the components of cermet compositions usu­ ally have different melting points. In sintering in the presence of a liquid phase, the sintering temperature lies between the melting point of the most fusible eutectic composition in the system and the melting point of the most refractory component. Thus, sintering is accomplished between the liquidus and solidus of the mixture, and the composition remains heterogeneous during the entire process.

-20- As long as the major part of the compact is not molten, the usual

ceramic or powder metallurgical methods can be employed.

Lenel stated further that sintering is most effective when the

refractory constituents are soluble to a limited extent in the liquid phase. If such solubility exists, the liquid will dissolve the small­

est particles of the refractory constituents and reprecipitate them

on the larger grains. This process will also occur during cooling if the phase which is liquid has little or no solid solubility for the refractory phases. By this process the grains of the solid phases will grow and voids in the compact will be eliminated. Voids are also filled when the surface tension of the molten phase pulls the grains together causing shrinkage of the compact.

Wertblad and Wulff^ found that if the molten phase does not dissolve the solid phase, then the liquid located in the grain boundaries hinders diffusion between the refractory particles and thus impedes sintering. Also, when no solubility between the components of a cermet composition exists and there is a sufficient amount of the low-melting phase in the body, the liquid will form a continuous

"matrix" which will hold the unmelted phases together on cooling.

The strength of such sintered compacts must be at least as great as 26 the strength of this matrix. Lenel thinks that dense bodies would result without solubility, if the molten phase completely wetted the solid particles.

-21- Grain Growth

Engle^ reported, that the particle sizes of the grains in cermet compositions are determined largely by the sintering treat­ ment. In commercial practice the time and temperature of sintering are chosen to avoid excessive grain growth. Fine-grained cemented carbides have grain sizes which are usually less than six microns in diameter, while a coarse-grained composition may contain grains as large as fifteen microns in diameter.

Skaupy^- found that, in general, the strength of cermets increased with decreasing particle size. Oswald-^ reported that hardness also increased with decreasing particle size. The decrease in hardness was not as pronounced as with common metal alloys, since the hard­ ness of the cermets was so great initially.

-22- MODE OF INVESTIGATION

The initial phase of this research was an attempt to develop a

cermet composition suitable for use in gas turbine blades from one

of the most oxidation resistant compositions found by Accountius^ in his investigation of the titanium carbide - silicon carbide - boron carbide system. The composition investigated, which consisted of 70$ TiC, 20$ SiC, and 10$ B,C by weight, was assigned the mnemonic 4 code number "721".

Sintering tests were made on the 721 mixture, and wetting tests performed to determine whether or not cobalt, nickel, chromium, and the intermetallic compound NiAl would form a bond with the 721 com­ position. Bonding was obtained between each of these metallic phases and the mixed carbide, after suitable heat treatment. Each of the metals and nickel-aluminum. were then mixed with the 721 composition in powder form, and these cermet compositions were formed into pellets.

The pellets were subjected to sintering tests to determine the temper­ atures at which the densest structures were obtained. Since the sintering tests on the 721 mixed carbide did not produce any strong specimens, and since no dense pellets were obtained from sintering tests on the cermet compositions, the investigation of the 721 mixed carbide was discontinued. Failure of these compositions to produce dense, strong bodies was attributed to the relatively large amounts of free graphite formed in reactions during sintering and to the presence of silicon carbide in the body.

-23- The second phase of this investigation was a study of a

material which contained titanium, boron, and silicon in the same

relative proportions as did 721. In the new composition, titanium

diboride replaced the boron carbide and part of the titanium carbide

of 721, and this substitution eliminated the reaction between the two carbides in which graphite was liberated. The silicon carbide

of 721 was replaced by silicon metal. The new composition consisted of 55«3$ TiC, 28.7^ TiBg* and. 16.($ Si by weight. It was hoped that this material would have the oxidation resistance of 721 but would have better sintered properties. Specimens of this composition alone and in combination with each of the metals cobalt, nickel, chromium, and the intermetallic compound NiAl were fired at various temperatures to determine whether or not dense specimens could be obtained.

The densest specimens produced in the second phase of the inves­ tigation were subjected to a preliminary oxidation test. A brief study was then made of the effect of reducing the titanium diboride and silicon contents of the most oxidation resistant cermet compo­ sition. Two cermet compositions were then selected for further evaluation in which bar-type specimens were used in determining modulus of rupture at room temperature. The composition with the highest room temperature strength was tested for modulus of rupture at elevated temperatures and for thermal shock resistance.

-24- MATERIALS AND EQUIPMENT

MATERIALS

Highly-Refractory Compounds

The source, particle size, and composition of each of the car­

bides and the titanium diboride used in this investigation are shown

in Table I. These materials were prepared for use by grinding them

in steel mills with steel balls. The titanium carbide, boron carbide,

and titanium diboride were milled for fifty hours, and the silicon

carbide was milled for seventy-two hours. Originally, methyl alcohol

was used as a milling vehicle to prevent oxidation of the newly

exposed surfaces of the grains. Gas pressures, which were large

enough to cause part of the mill charge to escape around the seal­

ing gasket, were formed during m3.lH.ng in the mills containing

titanium diboride and methyl alcohol. For this reason, carbon tetra­

chloride was then used as the milling medium, and later benzene was

used because a reaction between carbon tetrachloride and titanium

diboride and/or titanium carbide was encountered.

The titanium carbide, boron carbide, and silicon carbide were

purified to rid them of iron impurities picked up during milling.

The carbide to be purified was first leached with a lsl ratio by

volume of concentrated hydrochloric acid and distilled water. The material was then washed with distilled water until it was neutral

as tested by litmus paper. Finally, it was washed with methyl alcohol

-25- TABLE I

DATA ON HIGHLY-REFRACTORY MATERIALS

Partial Material Source Grade Analyses (50

Ti C Kennametal, Inc* -200 mesh Ti 80.0 Total C 20.0 Free C under 0.5

Si c Carborundum Co. FFF (—1,00 mesh) Si 69.3 _c 29.7

B.C Norton Co. #400 Flour B *T 78. C 20

TiB0 Norton Co. -20 mesh Ti 66.59 4 B 31.16 in a Buecnner filter and then air dried. The particle sizes of the milled carbides were determined by the Casagrande hydrometer method.

The particle size distributions obtained are given in Table II.

The first milled batch of titanium diboride was also purified in the above manner. However, during the process of washing the material with distilled water, evolution of a gas was observed; and, when the material was air dried, a small amount of a white crystalline substance appeared on the surface of the boride. Subse­ quent batches were milled with cobalt-bonded tungsten carbide slugs, and these batches were not purified.

Metals

In Figure 1, Page 11, the metals which form adherent protective oxide coatings are underlined. Of the underlined metals, only nickel, cobalt, chromium, and silicon were used in this investigation. Iron does not possess sufficient oxidation resistance, and the remainder of the underlined metals are either not sufficiently refractory or they are not available in large quantities. The silicon, nickel, cobalt, and electrolytic chromium used in this investigation were commercial grades. The intermetallic compound NiAl was prepared in the laboratories of The Ohio State University by Clinton C. McBride.

The source and approximate chemical compositions of these materials are shown in Table III. Silicon was prepared for use by milling the

-325 mesh metal for fifty hours. Cobalt was procured in two sizes; one was -325 mesh and the other was 98% between 2 and 4 microns

-27- TABLE II

PARTICLE SIZES OF MILLED* CARBIDES

______PERCENTAGE Particle Size TiC SiC ®4^ Rang®______'______

Greater than 10 microns 0. 3» 2.

5 to 10 microns 1 4 * 23. 26.

5 to 2 microns 82 . 52. 62.

Less than 2 microns 4* 22. 10.

^Titanium carbide and boron carbide were milled fifty hours. Silicon carbide was milled seventy-two hours.

-28- TABLE III

DATA ON RAW MATERIALS

Material Type Particle Size Partial Supplier Analysis

Chromium Electrolytic -325 meah Cr 99.3 Charles Hardy, inc.

Cobalt -325 mesh Co 98. Charles Hardy, Inc.

Cobalt 98$ 2-4 microns Co 98. Charles Hardy, Inc.

Nickel -325 mesh Ni 98. Charles Hardy, Inc.

NiAl Internetallic -325 mesh Ni 69.9 Prepared at The Ohio State Compound A1 Balance Univ. by C. C. McBride

Silicon -325 mesh Si 97.0 Charles Hardy, Inc. average diameter. The remainder of the metals were -325 mesh.

EQUIPMENT

Grinding Equipment

When a reduction in particle size was necessary, the material

was ground in a one-quart hardened manganese steel mill. A charge

of one-half inch steel balls or of cobalt bonded tungsten carbide

slugs with an average size of 1" x 0.75" x 0.5" was used. The mill­

ing vehicle was either methyl alcohol, carbon tetrachloride, or

benzene.

Forming Equipment

All of the cups used in the wetting tests and the pellets used

in the sintering tests were formed in small two-punch cylindrical

dies. Pressure was applied to the punches by use of a hydraulic jack. Bar specimens were formed by using a two-punch die in a large hydraulic press. The bars were approximately 4.5" x 0.5" x 0.2".

All of the pellets and bars were pressed hydrostatically after they were formed. This was accomplished by placing the specimens in special rubber envelopes which were then evacuated of air and sealed with rubber bands. The envelopes and their contents were immersed in a mixture of water and a water-soluble oil in a cylinder of specially treated tool steel. A piston was placed in the top of the cylinder, and a pressure of 500,000 pounds was applied to the piston,

-30- resulting in a pressure of approximately 35,000 p.s.i. on the liquid.

Vertical Molybdenum-Resistor Furnace.

A tube 1.75" in diameter, which was made from a sheet of molyb­

denum, served as the heating element and the firing chamber of the

vertical molybdenum-resistor furnace. Wetting test specimens were

placed on tungsten plates which were suspended in the firing chamber

by tungsten wires. This furnace could be evacuated to low pressures

because it contained a minimum amount of refractories to be out-

gassed. After evacuation, purified helium was allowed to flow into

the furnace, and this flowing helium atmosphere was maintained during

the firing cycle. The helium was purified by passing it through an

activated charcoal trap immersed in liquid air or liquid nitrogen.

The details of this furnace are shown in Figures 3 and 4*

High Frequency Induction Furnace

The sintering test specimens were fired in a large induction

furnace which was supplied with power by a 10,000 cycle 50 kilo­ watt motor-generator. This furnace was outstanding because of the low pressures and high temperatures which could be attained in the firing chamber. Sintering temperatures as high as 4500°F. have been obtained with this equipment. In this investigation, firings were made both in vacuum and in neutral atmospheres. Evacuation was accomplished with a mechanical vacuum pump and an oil diffusion pump. Details of this furnace are shown in Figures 5 and 6.

-31- Figure 3

Molybdenum-Resistor Furnace

-32- ATM OSPHERE' O C S T

ATMOSf>MfJERE IN

iR COPPER EXRAHSJp. NEOPRENE GASKET RING—

. BERYLUA R IN G S SREQ,

j h

TO OIL D IF F U S /O rV llllllttl\ PUMP

7i COOUN6 T t / B / N G S COPPER

TUNGSTEN R /3 D JA T /O N SHIELDSL

TUNGSTEN S H E E T S

ALUMINA SUPPORTS

COPPER EXPANSION R /A U G

£ TUBE

<5 RUBBER T O *= > £ * £ R

Figujrr© 4

Section Through Centero f Molybdenum?—Resistor Furnace

- 3 3 — Figure 5

High Frequency Induction Furnace

-34- Mice!

Figure 6

Section Through Center of High Frequency Induction Furnace

-35- Specimens were set on graphite discs which had been faced with titan­ ium carbide to prevent the specimens from becoming contaminated with carbon.

Globar Tube Furnace

A picture of the Globar tube furnace is shown in Figure 7. A silicon carbide tube heating element extends through the center of a metal case containing refractory brick, and water-cooled copper electrodes are attached to each end of the tube. A porcelain tube

1 .625” I» b. with one end tapered to a small opening extends through the silicon carbide tube and comprises the firing chamber. A uniform heat zone approximately 12” long exists in the chamber during firing.

Temperature is controlled by transformer settings and by an interval timer.

Hot Modulus of Rupture Furnace

The furnace in which modulus of rupture determinations were made at elevated temperatures is shown in Figure 8 , Specimens were sup­ ported on silicon carbide "knife-edges” and were loaded in the center of a three inch span by means of a silicon carbide rod connected to a .lever arm.

-36- Figure 7

Globar Tube Furnace Figure 8

Hot Modulus of Rupture Furnace

-38- PROCEDURE

TESTS FOR PHYSICAL PROPERTIES

The following tests were used to determine the physical proper­ ties of the specimens produced in this investigation.

Apparent Porosity

In the determination of the apparent porosities, each specimen was first weighed dry and then boiled for five hours in distilled water. The specimen remained in the water until it had cooled to room temperature, and then the weight of the specimen suspended in water was obtained. Its saturated weight was also determined after the excess water had been removed with a damp cloth. The density of water was assumed to be 1 gm./cc. Percent apparent porosity was calculated in the following manner:

Saturated weight - dry weight = volume of pores

Saturated weight - saturated suspended wt. = volume of specimen

% apparent porosity ■ — Q£„„P.oSgs— x 10O volume of specimens

Modulus Of Rupture

Determination of the modulus of rupture, or the maximum outer fiber stress, was made by supporting the specimen at two points and applying a uniformly increasing load at the center of the span until the specimen failed. The span was either 3«0" or 1.5" in all cases.

-39- Modulus of rupture was calculated by the following formula:

Modulus of rupture = J3., ^ 2 W d2

Where P is the load (in pounds) which caused failure

L is the length of the span in inches

W is the width of the specimen in inches

d is the depth of the specimen in inches.

The above formula is valid only for specimens with a rectangular cross section such as the bar specimens produced in this study.

Oxidation Resistance

Specimens to be tested for oxidation resistance were measured, weighed, and then placed in a Globar furnace which was maintained at

2000°F. The specimens were exposed to the oxygen atmosphere of the furnace. No attempt was made to seal the firing chamber, and cracks in the door permitted passage of air into the furnace. The specimens were removed from the furnace and weighed at various time intervals.

Weight gain per unit of surface area was calculated and used as a measure of oxidation. The parameter "k" as explained on page 12 was also calculated.

Thermal Shock Test

A specimen approximately 2” x 0.4” x 0.08" was held in the flame of a gas torch so that one edge was very near to or touching the inner cone of the flame. The bar was heated on the edge because greater thermal stresses were established in this way than could have been obtained by heating it on a face. When the temperature of the upper

-40- edge of the bar reached 2000°F., the specimen was quickly removed from the flame and quenched in an air stream for thirty seconds.

Cooling from 2000°F. to less than a red heat occurred in approxi­ mately two seconds. The flame was adjusted so that heating of the bar to the desired temperature required from twelve to fifteen seconds. One heating and one cooling operation made a cycle. After a specimen had been cycled a predetermined number of times, it was broken in a room temperature modulus of rupture test so that the effects of thermal shock could be studied.

TITANIUM CARBIDE - SILICON CARBIDE - BORON CARBIDE BASE COMPOSITIONS

Preparation of Compositions

Composition 721 (70$ TiC + 20$ SiC + 10$ B^C by weight) was pre­ pared by placing the correct ratios of the three carbides in a porce­ lain mill to which rubber stoppers were added to facilitate mixing.

The mill was then turned on the milling rack for fifty hours.

Four briquettes of 721 were dry pressed at 10,000 p.s.i. and fired in the large induction furnace. A firing schedule of about o _ 700 F. per hour was maintained. Near 3100°F. the sight tube became completely obstructed by material which apparently consisted of sulfide impurities in the graphite crucible and in the lampblack insulation. The maximum temperature reached could not be ascertained, but it was approximately 3900°F. A pressure of less than 350 microns was maintained in the furnace during the firing operation. The

-41- purpose of this firing was to react the 721 to form 721 R. In this

investigation, the suffix !'R" in a composition code number indicates

that the material was prereacted at elevated temperatures. X-ray

analysis revealed that the reaction to form titanium diboride was

complete.

The 721 R material was crushed by passing it through a jaw

crusher and a disc pulverizer until the material passed 35 mesh. It was then milled with steel balls for 100 hours in methyl alcohol.

The 721 R was air dried, purified, and screened through a 65 mesh

sieve. While the material was being washed with distilled water, a reaction occurred in which a gas was evolved. When the material was dried, a small amount of a white, crystalline material was deposited on the surface of the 721 R by the evaporating water.

This deposited substance appeared to be the same material that was found when titanium- diboride was purified, and it was identified microscopically as sassolite (B202*3H20). This reaction did not alter the composition of the 721 R from that of the same material prepared without milling or purification sufficiently to be detected by x-ray examination. Nevertheless, subsequent batches containing titanium diboride were milled with cobalt-bonded tungsten carbide slugs, and these batches were not purified.

Wetting Tests

Small circular tile, which were 0.75" in diameter and 0.25" thick, were formed by dry pressing 721 and 721 R. The 721 R tile

-A2- used in the experiments at 3050°F. and 3200°F. were prepared by form­ ing 721 tile which were then reacted or presintered at 3875°F. for one hour in an atmosphere of flowing argon. The top of each tile had a depression 0 .625” in diameter and 0 .125” deep which held the metal powder during the test. All of the tests were conducted in a neutral atmosphere. The metals tested and the test conditions are given in

Table V. The tests conducted at 3050°F. and 3200°F. were made in the vertical molybdenum-resistor furnace, and the vacuum induction furnace was used to fire the wetting test specimens at 3875°F.

The degree of wetting was determined by visual observation of a cross section of each specimen. A metal which infiltrated all of the pores of a tile was considered to have excellent wetting charac­ teristics under the conditions of the test. This was evidence of a high degree of wetting since the metal had to rise by capillary action to infiltrate the pores in the wall of the cup. Capillary action will not occur unless the liquid phase wets the solid phase. On the other hand, the formation of a bead of the metal on the mixed carbide tile was considered to indicate that the metal had very little tend­ ency to wet the refractory mixture under the conditions of the test.

Sintering Tests On Base Composition

Since it was desirable to determine the strength of the sintered

721 composition without a binder metal, bar specimens which could be tested for modulus of rupture were used in this portion of the investigation. Bars were formed from 721 and 721 R by the dry press-

-43- ing technique. After forming, the specimens were placed in special

rubber envelopes and pressed hydrostatically at 35,000 p.s.i. They were then set on a titanium carbide setting plate which was placed

inside a graphite crucible in the vacuum induction furnace. The maximum temperature attained and the furnace atmosphere for each

firing are given in Table VI. In every case, the ma-vimnm temperature

was maintained for one hour. Porosity and modulus of rupture at room

temperature were determined for each specimen.

Sintering Tests On Cermet Compositions

Sintering tests on cermet compositions based on 721-R were con­ ducted with pellets because this type of specimen required a minimum of material and produced the desired information concerning the optimum sintering temperature for each combination. Four cermet compositions were prepared by wet mixing 721-R with each of the following: chromium, cobalt, nickel, and the intermetallic compound

NiAl. The amount of metal added was equal to twenty percent of the weight of the mixed carbide. Thus, each composition consisted of

83.3$ 721-R + 16.7$ binder phase on a weight basis. The code desi­ gnations which were assigned to these and to all other compositions in this study were composed of two parts. The first group of numbers and letters indicated the base compositions, and the second group of letters indicated the binder metal. Thus, the symbol "721-R + NiAl" designated a body having as its base 721R composition with the inter­ metallic compound NiAl added as a binder metal. Compositions of all

-44- bodies tested in this investigation are given in Table IV.

Pellets 0.$" in diameter and approximately 0,3" thick were formed

at 25,000 p.s.i. by the dry press technique from each of the four com­

positions. Three specimens of each cermet material were sintered in

an atmosphere of purified helium at 100°F. temperature intervals in

the range 3300° to 4100°?. The sintering temperature was attained in

approximately one hour of firing, and a soaking time of one hour was

used in all cases. Porosity determinations were made on all sintered

pellets.

Bar specimens were prepared from composition 721 R + Go and

sintered in an atmosphere of purified helium for one hour at 150°F.

temperature intervals between the temperatures of 3600° and 41509F.

Room temperature modulus of rupture and apparent porosity were deter­ mined for each of these sintered bars.

TITANIUM CARBIDE - TITANIUM DIBORIDE - SILICON BASE COMPOSITIONS

Preliminary Evaluation

An investigation was made of bodies based on a composition con­

taining 55*3$ TiC 28,J% TIB2 + 16.0$ Si by weight. This combination

was designated by the numeral "III". The constituents for III compo­

sition were wet mixed in carbon tetrachloride, air dried, screened

through a 65 mesh sieve, and pressed into compacts which were pre­

sintered for one hour at 3600°F. in the large induction furnace in

a helium atmosphere. Cermet pellets were then prepared from these

compacts with the same metals and in the same manner as described

-45- TABLE XV

COMPOSITIONS STUDIED

Raw Materials - Weight Percent Base Composition______• Binder Metals Code Designations* TiC SiC B4C TiB2 Si Cr Co Ni NiAl

721 70.0 20.0 10.0 721 ♦ Cr 58.3 16.7 8.3 16.7 721 ♦ Co 58.3 16.7 8.3 16.7 721 * Ni 58.3 16.7 8.3 16.7 721 ♦ NiAl 58.3 16.7 8.3 16.7

III 55.3 28.7 16.0 III ♦ Cr 46.1 23.9 13.3 16.7 in ♦ Co 46.1 23.9 13.3 16.7 in ♦ Ni 46.1 23.9 13.3 16.7 ni + NiAl 46.1 23.9 13.3 16.7

III B ♦ Co 55.4 17.9 10.0 16.7 niCfCo 64.6 12.0 6.7 16.7 in d ♦ co 74.0 6.0 3.3 16.7

* The suffix nR", which indicates prereacted base compositions, has been omitted from these code designations. The compositions apply for both prereacted and unreacted base materials. for 721 R base compositions, and they were sintered for one hour in

helium atmospheres at 100°F. intervals in the temperature range of

2700° to 4100°F. Porosities of all specimens were determined, and

the densest specimens were subjected to oxidation tests.

Final Evaluation

The final phase of this investigation consisted of further evaluation of the most promising composition, III R + Co, found in the oxidation tests. The first step was the preparation of a series

of three cermet compositions which were the same as III R except that the amounts of silicon and titanium diboride were reduced by 25% intervals. The composition in which the silicon and titanium diboride contents were reduced by 25% from the amount in III R was designated as III R Bj the composition in which the original amounts of silicon and titanium diboride were reduced by 50% was called III R Cj and the composition which contained only one-fourth of the original amounts of these two materials was given the symbol III R D. Each of these base materials was wet mixed with 20% Co by weight, made into pellets, and sintered in a helium atmosphere at 3000°F. for one hour. The pellets were then tested for porosity and oxidation resistance.

On the basis of these oxidation tests, III R + Co and III R B

Co were selected for final evaluation. In addition, a batch of the same composition as III R B + Co but in which the base material had not been prereacted was prepared by milling the constituents in benzene for twenty-five hours. This composition was assigned the

-47- code symbol "III B * Go". The above three cermet compositions were

formed into bars which were sintered at temperatures between 2600°

and 3000°F. They were fired in a purified helium atmosphere in the

small Globar tube furnace. The sintering temperature was reached in about five hours of firing and was maintained for one hour. Porosity and modulus of rupture at room temperature were determined for all

sintered bars. Bars of composition III E B * Co were subjected to hot modulus of rupture tests at l600°, 1800°, and 2000°F«, to oxi­ dation tests at 2000°F., and to thermal shock tests. RESULTS AND DISCUSSION

TITANIUM CARBIDE - SILICON CARBIDE - BORON CARBIDE BASE COMPOSITIONS

Reactions During Preparation Of Raw Materials

TNheu titanium diboride was washed with distilled water in the

purification process, evolution of a gas was observed, and a white

crystalline material was deposited on the surface of the titanium

diboride during drying. This deposit was identified microscopically

as the mineral sassolite, B2O3 • 3H;>0. The following reaction was

assumed to have occurred between the titanium diboride and the water:

TiB2 + aH2° -> B^03 • 3H20 + Ti02 + 5H2 ^

The same reaction occurred when composition 721 R (70$ TiC + 20$ SiC

+ 10$ B/jC by weight) was purified because titanium diboride was present

as a reaction product. After the sassolite had been identified, all

future batches containing titanium diboride were not purified after

the milling operation. Cobalt-bonded tungsten carbide slugs were used instead of steel balls because the slugs introduced less con­

tamination into the batches because of their great hardness.

The reason for the high vapor or gas pressures built up in the mills during the grinding of batches containing titanium diboride in

methyl alcohol is not known. Since these pressures were great enough

to cause portions of the batches to be lost through blowouts in the

gaskets with which the mills were sealed, other milling liquids were

sought.

Carbon tetrachloride was the next liquid used as a suspending

-49- agent for milling batches containing titanium diboride. Excessive pressures were not formed in the mills when this liquid was used, and there were no immediate indications that any reaction was occur­ ring. However, when bodies made from the batches that had been milled in carbon tetrachloride were fired, a gray material deposited on the lid of the large induction furnace. This material was analyzed quali­ tatively and found to be composed almost entirely of titanium and chlorine. The crystals of the material could not be identified micro­ scopically because they were too small. Another indication that some reactions were occurring was that one of the batches of III R composi­ tion burned by spontaneous combustion as it was being air dried, after having been milled in carbon tetrachloride. This seemed unusual since carbon tetrachloride is noninflammable and is used as a fire extin­ guisher.

The above facts indicated that a reaction took place between the carbon tetrachloride and either the titanium carbide or the titanium diboride. Both of these reactions may have occurred. The following reaction was assumed to have occurred:

TiC + CC1. TiCl, + 20 k 4 Thermodynamic data^ were obtained, and calcualtion of the change in free energy for this reaction at 298°K. yielded a negative value in excess of ninety kilocalories. Thus, the assumed reaction is thermo­ dynamically plausible. Wulff^ milled titanium carbide cermet composi­ tions in carbon tetrachloride, and he did not report any such reaction.

He did find it necessary to presinter this material to eliminate "gas pockets", which may have resulted from the volatilization of titanium

-50- tetrachloride. If the assumed reaction does occur, the titanium te­

trachloride, which is a liquid at room temperatures, evidently de­

composes* %f on heating into one or more titanium chlorides which are solids at room- temperatures. Thermodynamic data were not available

for titanium diboride; and, therefore, calculations could not be made to determine whether or not reactions were possible between this

compound and carbon tetrachloride.

Wetting Tests

Results of the wetting tests are summarized in Table V. The wetting characteristics of 721 and 721 R were essentially the same.

The minor differences in the appearance of the specimens were attri­ buted to the differences in the rigidity of the tile. The portions of the 721 tile which were not Impregnated during firing were very weak, but the unimpregnated portions of the 721 R tale were strong because the tile had not been presintered.

During the tests at 3050°F. and 3200°F., the cobalt and NiAl melted; but they did not spread out on the surfaces of the tile or infiltrate them. These two metals were judged to have poor wetting * characteristics with 721 and 721 R at these temperatures. At 3875°F. the cobalt and NiAl wet the 721 R very well. A photomicrograph of the 721 R tile infiltrated by cobalt is shown in Figure 9, and the

721 R tile infiltrated with NiAl is shown in Figure 10. Comparison of these two photomicrographs reveals that grain growth was much greater and that more solution of the refractory phases occurred

-51- TABLE V

RESULTS OF WETTING TESTS

Mixed Carbide 721 721 721-R 721-R 721-R

Temperature 3050°F. 3200°F. 3050°F. 3200°F. 3875°F

Time at Temp* 0.5 Hr. 0.5 Hr. 0.5 Hr. 0.5 Hr. 0.5 Hr vi Atmosphere Helium, Helium Helium Helium Argon i

Metals

Co Poor Poor Good

Ni Fair Fair Good

CR Good Good Good

NiAl Poor Poor Poor Poor Good Figure 9

380 x Unetched. Composition 721 R + Co sintered at 3875°F. for one hour. Notice how much larger the grains are in this specimen than in 721 R + NiAl specimen shown in Figure 10.

-53- Figure 10

380 x Unetched. Composition 721 R + NiAl sintered at 3875°F. for one hour. Notice that the grain size is much smaller in this specimen than in 721 -*• Go specimen showi in Figure 9. when cobalt was used as the infiltrating metal than when NiAl was used.

The portion of the 721 R tile directly beneath the position formerly

occupied by the NiAl powder was porous, but the remainder of the im­

pregnated specimen was comparatively dense.

When nickel was tested at 3050° and 3200°F,, it appeared to have

wet the surfaces of both the 721 and 721 R tile but apparently did not

penetrate the tile. At 3875°F» nickel infiltrated the tile completely;

however, the grains of this impregnated tile were loosely bonded, and

they pulled out so easily that a polished section could not be made

from this specimen.

Chromium was the only metal vhich wet 721 and 721 R completely

at a temperature near the melting point of the metal. At 3200°F.

chromium infiltrated the tile and reacted with it to such an extent

that the impregnated specimen lost its original shape and assumed the

form of a button. This also happened at 3875°F., but this specimen was quite porous and weak. In these tests the wetting characteristics

of chromium are probably related to the pronounced tendency of this metal to form carbides.

Sintering Tests On Base Composition

Vacuum was used instead of an inert atmosphere in the first

attempts to sinter 721 because it was felt that sintering would be

facilitated by the removal of gases absorbed on the surfaces of the

carbide grains. Table VI shows that specimens sintered under a

vacuum of less than 350 microns were weak and porous.

-55- TABLE VI RESULTS OF SINTERING TESTS ON 721 BASE COMPOSITION

(Apparent Porosity in Percent and Modulus of Rupture at 75°F. in p.s.i.) Sintering Temperature and Atmosphere Composition 3400 (Vac.) 3600°F. (Argon) 3800°F. (Helium) 3875°F. (Argon) % p.s.i. % p.s.i. % p.s.i. % p.s.i. 721 33.2 * 25.3 7,200 15.8 5,600 18.8 12,200 33.2 # 27.8 * 14.2 8,000 18.2 11,400 33. 4 31.8 1 6. 4 11,500 34.6 * 30.3 1 6 . 4 7,100 34.3 * 9.7 Cracked Av. 33.7 28.8 14.5 8,000 18.5 11,800 721 R 10.6 9,900 15.6 21,800 11.6 11,100 15.7 7,500 Av. 11.1 10,500 15.6 14,600

Sintering Temperature and Atmosphere Composition 4000°F. (Vac.) 4000°F. (Helium)** 4100°F. (Argon) % p.s.i. % p.s.i. % p.s.i. 721 14.0 * 1 8 . 4 * 8.7 20,900 9.5 * 1 8 . 4 * 7.3 Cracked 9.4 * 32.0 * 9.9 4,800 13.8 * 14.2 * 18.8 * Av. 1 1.4 20.3 8.0 20,900 721 R 12.0 * 14.1 * 10.5 29,500 8.1 # 15.2 * 5.7 34,000 11.0 * 14.5 * 14.0 14.7 13.9 * 14.6 * Av. 11.8 14.6 8.1 31,800 * These specimens were very weak. They broke under the weight of the bucket in the M.O.R. test

** Fired to 3600°F. in vacuum and from 3600°F. to 4100°F. in a helium atmosphere When 721 and 721 R specimens were fired at 3600°F. in an atmo­ sphere of argon, no appreciable sintering occurred. Three subsequent tests at 3800°F., 3875°F«, and 4100°F. in atmospheres of either argon or purified helium produced sintered specimens which were still rela­ tively porous. The strength of these specimens increased and their apparent porosities decreased with increased firing temperature as is shown by Table VI. The strongest specimens produced in these sintering tests were made from 721 R and were fired at 4100°F. in an argon atmo­ sphere. They had an average modulus of rupture of 31,800 p.s.i. at room temperature and an average apparent porosity of 8.1$. A test was made at 4300°F. in an argon atmosphere, but the specimens fused to the titanium carbide setting plate so badly that they could not be removed.

In all cases in which the bars were strong enough for modulus of rupture determinations, the 721 R specimens had higher room temper­ ature strengths than the 721 bars. This difference was attributed to the more homogeneous structure of the 721 R specimens as compared to the 721 bars. A comparison of Figure 11, which is a photomicrograph of a 721 R specimen, and Figure 12, which is a photomicrograph of

721, shows the difference in structure between the two materials.

The areas of segregation are quite pronounced in the later photo­ micrograph, and they are almost completely eliminated in the 721 R structure. No such areas are shown in the photomicrograph of the

721 R, but a few very small areas of segregation are present in the specimen. Accountius^ postulated that these areas of segregation were formed during the reaction between titanium carbide and boron

-57- Figure 11

380 x Unetched. Composition 721 R sintered at 4100°F. for one hour. Notice the lack of areas of segregation. See Figure 12, which shows segregation in 721 composition sintered under the same conditions.

-58- Figure 12

380 x Unetched. Composition 721 sintered at 4100°F. for one hour. Dark gray phase is silicon carbide; light gray phase is titanium carbide; and the white phase is titanium diboride. Black areas represent voids or pits. Notice the apparent segregation of the material into titanium carbide rich areas and silicon carbide rich areas.

-59- carbide in which titanium diboride and graphite were produced. No reason for the apparent migration and concentration of the various constituents was advanced. It appeared that such a structure would more likely result from poor mixing technique, even though specimens produced from both dry mixed and wet mixed batches possessed the areas of segregation.

If poor mixing technique alone were responsible for the segre­ gation, then the greater homogeneity of 721 R bars as compared to

721 specimens would be produced by the milling of the 721 R material after it had been prereacted. Thus, milling a batch of unreacted 721 should also eliminate the areas of segregation in the sintered specimens.

A batch of 721 was milled for one hundred hours, and test bars of this material were sintered at 4100°F. for one hour in an atmo­ sphere of argon. The structures of these specimens could not be distinguished by microscopic examination from those of sintered specimens of 721 R. The new 721 specimens were slightly stronger than the 721 R bars sintered under identical conditions. This addi­ tional strength evidently resulted from the fact that the 721 material was milled fifty hours longer than the 721 R. Either the finer grain size or the additional cobalt picked up as milling contamination during the extra fifty hours of milling could have been responsible for the slightly higher moduli of rupture.

The desirability of mixing the constituents of cermet compositions by milling them together was made quite evident by the above experiment.

-60- Unfortunately, the raw materials used in this investigation were so expensive and so difficult to prepare for use that the size of indivi­ dual batches had to be held to a minimum. Equipment was not available for milling such small batches, which often consisted of less than one hundred grams of material. Therefore, the compositions for the pellets used in the sintering tests were prepared by wet mixing, but the re­ mainder of the bar specimens made in this study were prepared from compositions which had been milled for at least twenty-five hours.

Sintering Tests On Cermet Compositions

Results of the sintering tests on 721 R mixed carbide cermet compositions are shown in Table VII and Figure 13. None of the speci­ mens became dense during the firing operations. The 721 R base compo­ sition did not sinter dense without a metal addition, as reported earlier, but it was thought that a metal binder phase, which would be liquid at the sintering temperature, would facilitate sintering.

In general, cermets should have an apparent porosity of less than 0 .5$, if they are to exhibit their maximum strength and oxida­ tion resistance. This was the criterion established for bodies produced in this investigation. None of the 721 R +• Cr cermets attained a porosity of less than 10$ in the tests. Since chromium is known to have a great affinity for carbon, chromium carbides were probably produced in a reaction between the binder metal and the free carbon in the body. A comparatively large quantity of carbon was likely available from the reaction between titanium

—61— TABLE VII

APPARENT POROSITIES IN PERCENT OF SINTERED CERMET PELLETS BASED ON TIC 4 SiC 4 BjC

Sintering Temperatures

Compositions 3300°F. 3400°F. 3500°F. 3600°F. 3800°F. 3900°F. 4000°F. 4100°F.

7 2 1 R 4 Cr 18.2 16.8 16.3 13.1 18.0 24.9 21.4 20.9 17.5 18.4 15.8 13.3 17.2 25.1 33.9 17.5 20.3 15.0 10.9 17.1 29.6 26.2

Ave. 17.7 18.5 15.5 12. 4 17 . 4 24.9 25.3 27.0 721 R 4 Co 28.2 28.0 21.8 16.5 8.9 7.8 7.9 15.8 27.8 31.5 25.2 18.3 8.8 6 .4 7.8 17.7 25.0 32.9 22.2 22.1 9.8 6.2 12.7 16.8

Ave. 27.0 30.8 23.1 19.0 9.2 6.8 9.5 16. 8 721 R 4 N i 42.1 38.0 38.3 33.9 25.2 23.1 12.0 10. 4 41.8 39.4 38.8 34.8 26.4 22.7 11.1 9.0 41.6 39.7 38.6 34.6 26.1 23.6 10.1

Ave. 41.8 39.0 38.6 34 . 4 25.9 23.1 11.1 9.7

721 R 4 NiAl 32.5 31.1 30.7 25.8 14.0 15.7 4.5 27.5 32.4 30.6 30.3 27.1 15.1 16.2 6.0 32.6 30.8 31.0 28.3 16.1 15.0 4.5

Ave. 32.5 30.8 30.7 27.1 15.1 15.6 5.0 27.5 EFFECT OF SINTERING TEMPERATURE ON THE POROSITY Of fiOSiG+B+C BASE CERMETS

6 0-

6 0

i

zo

.1 »

0 0 3400 3soo 3600 37 00 3800 3900 4COO 4 1 0 0

WINTERING TEMPERATURE - cF. Figure 13

-63- carbide and boron carbide in which titanium diboride and carbon are

produced. If all of the chromium were converted into carbides in

this manner, then there would be no appreciable amount of liquid

present in the bodies at the sintering temperatures used in this

investigation; and sintering would occur essentially as a solid

state reaction.

The 721 R + Co cermets attained a minimum apparent porosity at

a sintering temperature of 3900°F. The average porosity of the spec­

imens at this temperature was 6 .8$, which is considerably greater

than the desired maximum apparent porosity of 0.5$. Each of these

cobalt-bonded pellets had on its surface many small areas which

appeared to be much more metallic than the remainder of the specimen,

and these areas contained relatively large holes. These areas of

high porosity were probably produced by poor distribution of the metal in the base composition and elimination of them should increase the density of the specimen considerably.

Therefore, a batch of 721 + Co was milled for twenty-five hours

in order to coat the ceramic grains with a thin metal layer. This material was then formed into bars which were subjected to sintering tests. Table VIII shows that bars were produced with less than two percent apparent porosity at sintering temperatures of 3900°F. and

4050°F. and that these specimens were still relatively weak. Higher sintering temperatures were not used because at 3900°F. and 4050°F.

substantial amounts of silicon carbide crystals were formed on the walls of the graphite crucible in which the bars were sintered.

-64- TABLE VIII

EFFECT OF SINTERING TEMPERATURE ON APPARENT POROSITY AND

MODULUS OF RUPTURE AT ROOM TEMPERATURE OF 721 R + Co

Sintering Temperature

3600°F ,’3750°F 39 0 0 °F 4050°F

M.O.R. Porosity M.O.R. Porosity Me 0

14,900 23.9 16,800 10.7 12,000 1.7 11,000 1 . 4

15,100 24.1 15,400 9.4 12,900 2.8 10,700 2.1

18,400 9.4 11,800 1.3

13,500 13.9 11,700 1.7

10,900 2.3

14,300 24.5 15,900 10.9 12,700 4.2 11,200 1.8

*This specimen had an unsintered core. This porosity value was not used in calculating the average The crystals were formed either by sublimation and recrystallization of the silicon carbide from the specimens or by decomposition of the silicon carbide in the 721 composition followed by volatilization of the silicon and reaction between this vapor and the graphite crucible to form the silicon carbide crystals. In either case, the loss of such large quantities of silicon from the cermet composition would lower its oxidation resistance appreciably. Furthermore, it is quite likely that a dense structure cannot be obtained at sintering temper­ atures above the sublimation temperature of silicon carbide, 3632°F.^, until all of the silicon carbide has been eliminated. Another reason for not sintering the 721 R base cermets at temperatures higher than

A100°F. was that the base composition alone fused when sintered at

4300°F., and the cermets would probably become fused masses at lower temperatures.

The fonowing phases were found by x-ray examination in a

721 + Co bar fired at 3900°F.: titanium carbide, titanium diboride, cobalt monosilicide, graphite, and one or more unidentified phases.

Titanium carbide was the only one of the original constituentj which could be identified in the sintered body. Both the reaction between titanium carbide and boron carbide and that between cobalt and silicon carbide should produce graphite as reaction products. This large amount of carbon in the body was probably the factor responsible for the low strength values of the 721 + Co cermets. Carbon contents as low as two or three percent by weight are thought to be extremely detrimental to the strength of titanium carbide base cermets. For

—66“ this reason Kennametal, Inc. produces titanium carbide by the menstrum process so that it will have a carbon content of less than one-half

of one percent.

Pellets of 721 R + Ni became denser as the sintering temperature was increased, and at 4100°F. they obtained their minimum average

apparent porosity which was 9.7%. These pellets were not sintered at higher temperatures for the reasons given previously.

Cermets of the 721 R + NiAl composition had porosities which decreased as the sintering temperature was increased up to 4000°F., yet the pellets were again quite porous when sintered at 4100°F.

This large increase in porosity may be attributed to either decom­ position of the silicon carbide with subsequent volatilization of silicon or to sublimation of the intermetallic compound NiAl. The densest specimens obtained had an apparent porosity of 4*5%. Micro­ scopic examination of one of these specimens revealed that it had small particle size, although it had been sintered at 4000°F. The silicon carbide had not reacted with any of the other constituents of the body to any appreciable extent.

TITANIUM CARBIDE - TITANIUM DIBORIDE - SILICON BASE COMPOSITIONS

Preliminary Evaluation

The failure of 721 base cermets to become dense when sintered under the conditions of these tests was attributed to the amount of free graphite in the bodies and to the presence of silicon carbide.

-67- Most of the free graphite was present as a product of the reactions which occurred during sintering. The main disadvantages of silicon carbide are that it is not easily wetted by the metals tested and that dense silicon carbide bodies can be produced only by elaborate techniques.

A new composition was sought which would not have free graphite produced as a reaction product during sintering but would retain the oxidation characteristics of 721 composition. The second composition studied contained titanium carbide, titanium diboride, and silicon, and it was designated by the code number "III". This composition contained titanium, boron, and silicon in the same relative propor­ tions as did 721. By replacing all of the boron carbide and part of the titanium carbide of 721 with titanium diboride, the carbon liber­ ated in the reaction between the two carbides was eliminated. Replac­ ing the silicon carbide of 721 with silicon metal should eliminate the free carbon produced in reactions between the binder metals and silicon carbide in which the silicides of the metals were formed.

Removal of the silicon carbide, which is not easily wetted or sin­ tered, should also be beneficial to the fired properties of the cermets.

Base composition III R was subjected to sintering tests in order to secure a dense specimen for the study of the effect of the addition of a binder metal on oxidation resistance. The results of all sintering tests made on pellets containing composition III R are given in Table IX.

-68- TABLE IX

APPARENT POROSITIES IN PERCENT OF SINTERED PELLETS BASED ON TiC 4 TiB2 + Si

Sintering Temperatures Compositions 2700°F. 2800°F. 2900°F. 3000°F. 3100°F. 3200°F. 3300°F. 3400°F.

III R 11.7 8.8 6.2 6.3 13.6 10.6 10.2 6.5 9.3 1 1. 4 10.0 9.7 6.5 8.4 7.9

Ave. 10.8 9.6 6.4 8.0 11.0 Ill R + Cr 8.5 5.9 2.8 8.4 13.0 30.0 47.1 40.4 9.0 6.7 5.3 6.8 13.6 31.2 50.7 41.1 8.9 7 . 2 5.3 6.2 15.6 32.7 49.7 42.4

Ave. 8.8 6.6 4*4 7.1 14.1 31.3 49.2 41.3 Ill R + Co 0.0 0.2 0.3 0.4 6.2 5.9 0.1 0.0 0.3 0.2 7.0 7.9 0.0 0.3 0.6 0.2 9.7

Ave. 0.0 0.2 0 . 4 0.3 6.6 7.8 H I R ♦ Ni 11.8 9.5 13.0 5.7 14.8 13.3 8.4 15.6 2.5 12.1 13.7 7 .8 15.7 6.8 11.7

Ave. 12.9 8.6 14.8 5.0 12.9 Ill R + NiAl 22.0 21.6 19.7 15.7 16.2 20.7 26.8 17.8 14.0 13.3 21.3 20.9 16.7 13.7 13.3

Ave. 21.3 23.1 18.1 14.5 14.3 TABLE IX (Continued)

Sintering Temperatures

Compositions 3500°F. 3600°F. 3800°F. 3900°F. A000°F. 4100°F. 3150°F.* 3400°F.* III R 18.6 14.5 13.0 25.9 3.0 11.0 16.7 1.8 15.1 6.1 16.0 24.2 8.4 7 . 4 20.1 11.5 8.5 . 18.7 2 5.8 11.8 16.0

Ave. 15.1 9.7 15.9 25.3 7.7 11.5 18. 4 1.8 Ill R ♦ Cr 50.0 40.8 8.8 46.7 47.2 51.5 43.3

Ave. 49.4 43.8 8.8 Ill R * Go 16.8 33.1 13.2 39.8 12.9 13.5

Ave. 14 . 4 33.1 13.2 39.8 Ill R ♦ Ni 15.6 27.1 21.8 U . 9 14 . 4 27.8 13.2 25.3

Ave. 14.4 26.7 21.8 11.9 Ill R ♦ NiAl 14.3 14.0 12.4 25.6 12.2 19.2 12.1 17.9 18.9 14.2 23.0 21.5 14*4 22.0 18.5

Ave. 16.1 15.8 13.3 23.5 17 . 4 19.2 12.1

*Sintered in vacuum. Other firings made in helium atmospheres. All of the pellets of composition III R sintered in a helium

atmosphere were relatively porous. None of the tests produced spec­

imens with an average porosity of less than six percent. However,

one specimen sintered in a vacuum at 3400°F. had a porosity of 1.8%; and, since this was the densest pellet of the basic composition, it was subjected to oxidation tests.

Actually, the basic composition III was a cermet material, since it contained ceramic phases and a metal phase. It is quite possible that this composition could be sintered dense in a vacuum at other temperatures, but a cermet with silicon as the binder metal would probably be too brittle for practical purposes.

Figure 14 shows that III R + Cr cermets had a minimum average porosity of 4«A% when sintered at 2900°F. These specimens had sharp outlines and were quite metallic in appearance. If a technique could be found for sintering these specimens dense, it is possible that they would possess unusually good high temperature properties.

Cermets of III R Ni composition reached a minimum average o porosity of 5.0% at a sintering temperature of 3300 F. None of the pellets were considered to be sufficiently dense for comparative oxidation tests.

Under the conditions of these tests, the intermetallic compound

NiAl was the poorest bonding metal tested. No specimens were obtained with a porosity of less than twelve percent..

Cobalt was the only binder metal tested with composition III R that produced dense pellets under the conditions of these tests.

-71- POROSITY 40 20 0 5 30 60 2000 OFTiC+TiBi*Si BASE CERMETS BASE OFTiC+TiBi*Si EPRTR O H POROSITY THE ON TEMPERATURE FET F SINTERING OF EFFECT 0 0 0 3 NTERI EMPERATURE - E R U T A R E P M TE G IN R E T IN S 0 0 2 3 Figure 14 Figure -72- EMPERATURE R U T A R E P M TE

0 0 8 3 °F. 000 4 Specimens with less than 0 .1$ porosity were obtained over the tem­

perature range 2900° to 3200°F. with cobalt as the binding phase.

The cobalt used had a grain size of 98$ between two and four microns,

and it was felt that this may have been the factor which was respon­

sible for the dense specimens obtained with this metal when other

binder metals produced porous bodies. Therefore, nickel with a grain

size of four microns was procured and used as a binder metal in

combination with composition III R. When sintered, specimens of

this material were not materially different from those in which the

binder metal was -325 mesh nickel. Therefore, it was concluded that

the grain size of the cobalt was not necessarily responsible for

the fact that these pellets were denser than those bonded with other metals. A photomicrograph of a pellet of composition III R bonded with cobalt is shown in Figure 15. This specimen was hard, and it

had a very fine-grained structure and an apparent porosity of zero.

Oxidation tests were made on specimens of III R +• Go and on

one pellet of composition III R. Results of the oxidation tests

are summarized in Table X. The values of the constant "k", the

significance of which has been explained on page 12, are given in

Table XI. Actually, the values of "k" varied widely during the

three hundred hour test. This indicates that the oxide coatings were not protective during the entire test. Figure 16 shows that

both III R and III R ♦ Co formed protective glass coatings during the first one hundred hours of oxidation^ therefore, "k" values for

both materials were practically constant during this time except for

-73- Figure 15

1500 x Unetched* Composition III R + Co sintered at 2S00°F. for one hour. This specimen has a very fine-grained struc­ ture and an apparent porosity of zero.

-74- TABLE X RESULTS OF OXIDATION TESTS

Composition Total Weight Gain in Units of 0.1 mg/cm2

10 25 50 100 200 300 Hours Hours H0UT8 Hours Hours Hours

III R 47.1 57.2 76.3 97.5 220. 239. i n r ♦ Co 11.8 17.7 29.0 46.1 104. 225. 11.3 18.6 30.7 46.7 165. 310. 12.1 118. 265.

Ave. 12.0 18.2 29.8 46.4 129. 266.

Ill RB ♦ Co 51.6 78.7 112. 148. 205. 244. 64.5 84.3 117. 148. 208. 252. 57.4 7 5 . 0 110. 138. 195. 236.

Ave. 57.8 79.3 113. 145. 203. 243.

Ill RC ♦ Co 63.8 88.8 141. 257. 535. 749. 70.5 95.3 144. 240. 495. 703. 72.7 99.2 153.

Ave. 69.0 9 4 . 4 146. 248. 515. 726.

Ill RD + Co 85*2 162 . 343. 82.8 150. 77 . 5 149.

Ave. 81.8 154. 343. • • • • • • • • • TABLE XI

VALUES FOR THE PARAMETER ''k1' FROM LONG TIME OXIDATION TESTS -4" cv' i $ 1 Parameter "k" •

Sample 10 25 50 100 200 300 Number Hours Hours Hours Hours Hours Hours

III R .0370 .0218 .0194 .0158 .0403 .0318

III R ♦ 136 .0024 .0022 .0030 .0036 .0139 .0393

III RB ♦ Co .0557 .04L9 .0425 .0350 .0344 .0328

III RC ♦ Co .0793 .0594 .0711 .1025 .221 .293

III RD + Co .111 .157 .393 80, 000

7 0 }0 0 0 Lons Time Oxidation In Open Air At ZOOO°f\

£>0,000

0 50 100 ZOO 300 TIME - HOURS Figure 16

-77- the initial value for III R. The initial oxidation resistance of

III R composition was improved by the addition of the metal. This

effect may have been created by the greater density of the cobalt

bonded cermets as compared to sintered III R material without a

binder metal.

During the one hundred to two hundred hour interval, the III R

composition oxidized at a very rapid rate, but the rate of oxidation during the following one hundred hours was somewhat less. This

suggests that the protective glass coating developed during the first hundred hours of oxidation was disrupted so that it no longer formed a protective coating.

Composition III R + Co also began to oxidize quite rapidly sifter the first hundred hours of oxidation. Figure 16 shows that, during the first hundred hours of the test, both compositions III R and

III R Co had much better resistance to oxidation than did the best commercial cemented csirbide, Kennametal's composition K-138 A.

The oxidation tests on K-138 A were made by Accountius^, who found that this cemented csirbide followed the parabolic law of oxidation} thus, during oxidation K-138 A forms a protective coating which inhibits further oxidation. At the end of three hundred hours of oxidation the three compositions had gained practically the same amount of weight per unit of surface area.

Not much importance can be attached to such comparisons as these because the weight gain per unit of surface area does not give a time picture of how much a body is oxidized. For example,

-78- Kennametal1 s K-138 A is composed mainly of titanium carbide and

similar carbides. When titanium carbide oxidizes the products of

oxidation are usually titanium dioxide and either carbon monoxide or

carbon dioxide. In either case, for each molecular weight of tita­

nium carbide oxidized one molecular weight of oxygen is added to

the weight of the body and one atomic weight of carbon is lost through

the formation of a gas. On the other hand, if part of the carbides

in Kennametal's K-138 A were replaced by titanium diboride and silicon,

a composition very similar to III R ♦ Co could be obtained. However, as titanium diboride oxidizes, titanium dioxide and boric oxide are

formed with the addition of two and one-half molecular weights of oxygen per mole of titanium diboride oxidized.

The above example serves to illustrate the futility of using weight gain per unit of area as a measure of oxidation resistance.

A reliable method.of measuring oxidation would probably involve the specific volume of the constituents of the cermets and of their oxides, the rates of the various oxide-forming reactions, and the properties of the oxides, as well as the other considerations given.

Any good method devised for measuring the effects of oxidation will likely apply only to one or, at most, several of the physical proper­ ties of cermets.

The weight gain per unit surface area oxidation test method is used because it provides a convenient method of making crude compar­ isons of oxidation resistance. If a reliable method were known for making these tests, it would be of little value until the oxygen pres­ sures in gas turbines have been determined. At present, the best

-79- method of comparing the oxidation resistance of various cermets which are to be used in jet engines is to prepare gas turbine blades from them, to subject them to actual operating conditions, and to observe the resultant deterioration. It is often difficult to determine whether such deterioration is caused by lack of oxidation resistance or by other factors.

Alteration of Cobalt-Silicon Ratio

Cobalt silicides were found in reacted III R ♦ Co pellets by x-ray examination, and it was assumed that all or nearly all of the silicon and cobalt in' this composition had reacted to form cobalt silicides. The only other phases which were present in sufficient quantities to be detected by this method were titanium carbide and titanium diboride. If only the cobalt and silicon present in

III R + Co cermets are considered, the composition of the system falls at 4 4 .4$ silicon on the cobalt-silicon phase diagram shown in

Figure 2, page 16. Under conditions in which equilibrium was attained, this system would consist of 26.7$ CoSi and 73•3$ CoSi2 crystals$ however, cobalt disilicide has been reported by Lebeau^ to be brittle.

Thus, a change in the cobalt-silicon ratio which would produce more of the cobalt monosilicide seemed desirable. The high temperature strength of the III R + Co body could probably be improved consider­ ably by altering the composition so that the only cobalt silicide formed (assuming complete reaction) would be the pure compound CoSi

(32.2$ silicon by weight). A slightly higher percentage of silicon

-80- would likely be even better, since in the case of incomplete reaction

such a composition would lessen the chances of forming cobalt sili­ cides which would fall into the range of compositions between Co2Si and CoSi, where there are solid phase transitions. Such transitions are not desirable in cermets.

Three obvious ways in which this alteration could be accomplished were: (l) to increase the relative amount of cobaltj (2) to add all or a portion of the silicon in the form of a compound which is more stable than the cobalt silicides over the temperature range of proc­ essing and usej and (3) to reduce the relative amount of silicon in the body.

Increasing the amount of cobalt did not seem practical, since thirty percent by weight of composition III R + Co was composed of silicon and cobalt. Commercial cermets usually have a maximum of thirty percent binder phase unless a relatively refractory metal such as chromium is used.

Titanium disilieide was considered as a material which might be more stable than the cobalt silicides at the temperatures encountered in the production and use of these cermets. The selection of tita­ nium disilieide was based on the fact that it was thought to be a refractory material which would not lower the refractoriness of the

III R Co composition unless it reacted with the other constituents of the body. Another reason for its selection was that use of tita­ nium disilieide would not introduce any new elements into the cermets.

VJhen pellets containing a ratio of cobalt to titanium disilieide of 2:1 on a molecular weight basis were sintered at 2700°F., the

two materials reacted to form cobalt monosilicide and unidentified

compounds. No cobalt, titanium disilieide, or titanium could be

found by x-ray examination. The unidentified phases were probably

ternary compounds. Since titanium disilieide reacted with the cobalt,

it was not investigated further.

Decreasing the relative amount of silicon in III R + Co was

attempted vfith some success. The titanium diboride content was

reduced at the same time that the silicon content was decreased so

that the ratio of silica to boric oxide would be constant in the

protective films formed during oxidation. Pellets of each of the

three compositions which had their silicon and titanium diboride con­ tents decreased by successive 25% intervals were sintered to less than one percent porosity at 3000°F. These pellets were then sub­

jected to oxidation tests. The results of these tests are shown in

Table X, and in Figure 17.

When the amounts of silicon and titanium diboride were reduced by one-fourth intervals, the initial rate of oxidation of the sin­ tered specimens increased rapidly. This loss of oxidation resistance was quite large for the III R B + Co cermets, which contained three- fourths of the original amounts of silicon and titanium diboride.

Figure 16, page 77# shows that III R B + Co obeyed the parabolic law of oxidation during the period of oxidation from fifty hours to three hundred hours. In this time interval the "k" values for III R B + Co averaged 0.0341 mg.^/cm. Vroin* £ 0.0013. On the basis of the oxida-

-82- 1

EFFECT OF CHANGES IN COMPOSITION! ON! O XID A TIO N OF TiC ■^TiB^.+Si + Co CERMETS 700

600

5 500-

? 4 0 0

O 200

100

5 0 IOO 150 200 250 HOURS OF OXIDATION

Figure 17

-83- tion test used, III R B ♦ Co compares favorably with Kennametal's

K-138 A composition, which had an average "kM value of 0*0386

mg.2/cm. fymin. ♦ 0.0016. The III R B + Co pellets had smooth sur­

faces after oxidation, but the surfaces of all of the other pellets

seemed rougher after oxidation than before the test.

Sintered composition III R D ♦ Co, which contained only one-

fourth of the original amounts of silicon and titanium diboride,

oxidized at a rate which indicated that the oxide coating formed on

these pellets had little or no tendency to inhibit further oxidation.

Figure 17 reveals that the weight gain per unit of surface area for

III R D * Co pellets plotted against time produced a curve which was almost linear. When these specimens were removed for weighing after one hundred hours of oxidation, all of the III R D + Co pellets adhered to the refractory brick on which they rested during the tests. The pellets were removed, and tests on them were discontinued because small grains of the refractory were imbedded in the glassy surface coating.

Pellets of III R C + Co, which contained one-half of the orig­ inal amounts of silicon and titanium diboride oxidized at a rate that was also linear in nature after the first ten hours. After three hundred hours of testing, the pellets of this composition had gained more than twice as much weight per unit of surface area than any of the other compositions with higher silicon and titanium diboride contents.

-84- Final Evaluation

Results of the preliminary tests indicated that only two of the

compositions investigated, III R ♦ Co and III R B + Co possessed

good sintering characteristics and good resistance to oxidation.

Therefore, bars of these two compositions were sintered at various

temperatures, and porosity determinations and room temperature modulus of rupture tests were made on the sintered specimens to

determine the optimum sintering temperatures for each composition.

In addition, the same tests were conducted on a composition desig­

nated as III B + Co. This material had the same composition as

III R B + Co, but the base composition of the former had not been prereacted, and it was milled with the binder metal for only twenty-five hours.

The results of these sintering tests are shown in Figure 18 and

Table XII. The strength of all three cermet compositions increased rapidly to a maximum and then decreased less rapidly as the firing temperatures were increased. The refractory components of these compositions were only slightly wetted by the liquid phases at tem­ peratures below 2600°F., and there was no evidence of the spreading type of wetting in this temperature range. Above 2600°F. the liquid phase wetted the base materials to a high degree, and the strength of the specimens increased rapidly. As the optimum sintering temperatures were exceeded, the bars lost strength, probably as a result of grain growth.

The fact that the spreading type of wetting did not occur at

-85- 6*4 Por. 3.1 3.0 4.8 9.7

3000°F. (p.s.i.) ($) 23,700 30,200 M.O.R. 14,900 3.9 0.9 1.2 0.8 18,000 0.3 0.7 0.8 0.3 1.0 0.6 0.5 0.6 34,900 1.8 0.6 29,400 0.5 0.6 0.8 • Por. .(*) 2900°F 51,800 56,000 (p.s.i.) 52,900 58,500 51,200 0.6 45,600 49,800 49,400 46,600 47,000 43,400 35,20064,300 0.2 44,900 40,900 47,600 44,300 37,400 M.O.R. : (*) 4.5 0.3 0.3 0.4 0.4 8.4 4*4 0.5 0.9 0.7 0.9 0.4 0.7 0.7 0.3 • Por. 2800°F (p.s.i.) 53,500 51,400 53,100 55,500 63,800 62,100 63,200 60,400 62,000 40,500 48,600 48,300 72,300 M.O.R. i TABLEXII (%) 0.4 0.5 0.5 1.9 Por. 10.0 13.7 13.1 12.4 12.7 16.9 17.1 2700°F. 9,200 (p.s.i.) 69,500 0.663,700 55,900 65,600 0.8 26,100 72,800 63,300 40,800 12.8 53,300 16,600 10,700 46,500 Sinteriqg Temperature 9,100 25.4 (p.s.i.) 13,700 25.8 60,800 13,200 24.3 M.O.R. Por. M.O.R. 18,600* 23.9 11,600 22.0 63,800 1.0 RUPTURE AT ROOMTEMPERATURE OF TiC + TiB2 + Si + Co CERMETS EFFECT OF SINTERING TEMPERATURE ON APPARENT POROSITYAND MODULUS OF Ave. Ave. 12,20015.6 Ave. 31,600 Composition 2600°F. Ill B 4- IllCo B 4- Ill RB + Co III R* Co 13,000 * Two specimens with less than7*000 p.s.i. were not included*

- 86- MODULUS OF RUPTURE - P. S. 70,000 50000 60,000 IOOOO EPRTR FRTCTB+Ho CERMETS FORTiC+T«Bz+SHCo TEMPERATURE T C E F F E N RNVRE TEGH T ROOM AT STRENGTH TRANSVERSE ON OF NT NG TEMPERATURE-°F. G IN R TE IN S NT NG G IN R TE IN S FIGURE FIGURE 2700 -87- TEMPERATURE 20 000 30 '2800 IQ 2900

lower temperatures was demonstrated by tests made at sintering tem­

peratures of 2500°F. and 2600°F. in the large vacuum induction fur­

nace. Figure 19 shows some of these specimens in the graphite boats

in which they were fired. The bottoms of these boats were covered with powdered titanium carbide, and broken pieces of titanium carbide bars were used to separate the specimens. Formation of metal beads

such as those shown on these specimens usually indicates that the refractory phase or phases have consolidated during sintering and forced the metal out. However, in this case the bars were still quite porous, and so the metal obviously was not forced out by con­ traction of the refractory skeleton. This metal was evidently forced to the surfaces by escaping gasesj and, since the temperature was not high enough for the spreading type of wetting reaction to occur, the metal had little tendency to go back into the pores or to spread out on the surfaces. Two of the specimens which are shown in

Figure 19 were fired again to a temperature of 2700°F. The metallic phases then spread over the surfaces of the specimens and infiltrated the pores. The higher temperature must have lowered the surface tension of the metal sufficiently to permit the spreading type of wetting to occur.

All of the specimens shown in Figure 19 contained titanium car­ bide and titanium diboride and were milled in carbon tetrachloride.

Decomposition or volatilization of the titanium chloride compound or compounds formed as a result of the milling operation, probably pro­ vided the gas pressure which forced the metal out of the pores.

-88- Figure 19

Graphite boats containing bars fired at 2600°F. Bars were covered with beads of metal, which indicates that the spreading type of wetting reaction had not occurred under the test conditions. All but one of the specimens in the first boat had disintegrated. Heavy deposits of the material which contained titanium and chlorine as its major constituents were formed on the lid of the furnace dur­ ing these two firings. All but one of the specimens in the first boat were completely disrupted. This was attributed to formation of high gas pressures in the closed poresj and, also, the firing cycle was probably much too fast. All subsequent firings were made in the small GLobar tube furnace with a longer heating time.

The strongest specimens were secured from composition III R B *s*

Co at a sintering temperature of 2700°F. Tiieir average modulus of rupture at room temperature was 65,600 p.s.i. Additional specimens of III R B + Co were sintered and found to be as strong as the orig­ inal bars; therefore, these results are reproducible. A photomicro­ graph of a sintered III R B + Co bar is shown in Figure 20. This specimen was sintered at 2800°F., yet it still has a fine-grained structure. It had an apparent porosity of 0.5% and a specific gravity of approximately 5.40.

The maximum average strength of composition III R Co was

51,400 p.s.i., which was attained at a sintering temperature of

2800°F. The cermets of III B + Co also had their maximum average strength of 53,100 p.s.i. when fired at 2800°F. The higher sinter­ ing temperature required and the lower strength values obtained for III B + Co specimens as compared to cermets of III R B t Co are probably the result of the difference in grain size. The

III R B + Co material was milled twenty-five hours longer than the III B ♦ Co.

-90- Figure 20

1500 X Etched. Composition III R B + Co sintered at 2800°F. for one hour. This cermet has a fine-grained structure.

-91- Since specimens of III R B + Co had the highest modulus of rupture at room temperture, this composition was selected for modulus of rup­ ture tests at 1600°F., 1800°F., and 2000°F•, and for thermal shock resistance tests. The results of the modulus of rupture tests at elevated temperatures are given in Table XIII and in Figure 21. The modulus of rupture of this composition was considerably higher at

1600 F. and at 1800°F. than at room temperature, and at 2000°F. the modulus of rupture was almost equal to the room temperature values.

During the tests the lever arm of the modulus of rupture fur­ nace was observed to move slightly after loading had been in progress for approximately twenty seconds. This indicated that the specimens were deforming to some extent. This did not appear to be plastic de­ formation since the lever arm did not move continuously, but moved a little and then ceased to move even though the load was being applied continuously® An. estimated 30,OCX) p.s.i. of outer fiber stress was applied to all of the specimens before any indications of deformation could be detected. The ability of this material to deform at elevated temperatures without rupturing probably reduced stress concentrations and was directly responsible for the increased strength. A maximum average modulus of rupture of 87,900 p.s.i. was attained at 1800°F,

Table XIV gives the results of the thermal shock tests. All of the specimens retained at least seventy-five percent of their orig­ inal strength in modulus of rupture tests after thermal cycling. The number of cycles to which the specimens were subjected after the first ten cycles apparently did not affect the strength of the specimens.

-92- TABLE XIII

CHANGE OF MODULUS OF RUPTURE OF COMPOSITION III RB •* Co WITH TEMPERATURE

MODULUS OF RUPTURE - (p.s.i.)

75°F 160Q°F 18QO°F 2000°F

63,300 81,700 89,700 64,800

69,500 84,900 86,100 66,800

60,800

63,800

63,700

72,800

Ave. 65,600 83,300 87,900 65,800

-93- MODULUS OF RUPTURE - P. S. 90,000 70,000 75 HNE F UU O RUPTURE OF E R U T P U R OF DULUS O M OF CHANGE OSTON O ILBC WI H IT W IELRB+Co OF N SITIO PO M O C E ~ R °F. U T A R E P M E T e r u t a r e p m e t 1600 P IUE 2/ 2 FIGURE. -94- 1800 2000 TABLE XIV

RESULTS OF THERMAL SHOCK TESTS

Number of Cycles Modulus of Rupture At Room Temperature After Cycling (p.s.i.)

10 50,400 10 41,100

Av. $0,800

20 48,100 20 57,200

Av. 52,600

30 57,600 30 48,300

Av. 53,300

¥> 52,600 40 52,800

Av. 52,700

50 62,000 50 51,400

Av. 56,700

-95- The thermal shock tests were conducted on specimens which were approximately 2" x 0.4" x 0.08". Whether larger specimens would have such good thermal shock resistance is not definitely known.

However, the thermal conductivity of these specimens appeared to be quite high, and it is thought that larger specimens would with­ stand thermal cycling without serious deterioration.

These laboratory tests have shown that composition III R B + Co possesses many of the properties desired in a material to be used in gas turbine blades. Tests should be made under actual service con­ ditions to determine definitely the value of this cermet composition for use in rotor blades. GENERAL SUMMARY

The purpose of this investigation was to develop a self-glazing

titanium carbide base cermet composition which would be suitable for use in gas turbine blades. Binder metals were combined with an oxi­ dation resistant material composed of 70$ TiC ♦ 20% SiC + 10% B^C by weight, and these cermet compositions were fired over a range of temperatures in an effort to produce dense bodies. Failure of these materials to sinter dense under the conditions of these tests elim­ inated them from further consideration.

The second phase of this study was an investigation of a base composition of the same materials as were present in the sintered mixed carbide, except that the amount of free graphite was greatly reduced and the silicon carbide was replaced by silicon metal. When this basic material was sintered with cobalt as a binder metal, dense, oxidation resistant cermets were obtained. Thus, either the large amount of free graphite or the presence of silicon carbide in the body or both of these factors were responsible for the failure of the original mixed carbide cermets to sinter dense.

An effort was made to adjust the silicon to cobalt ratio in the composition that had sintered dense so that the high temperature strength would be increased and brittleness would be reduced. This was accomplished without seriously impairing the oxidation resistance of the material. The resultant composition was 55*4# TiC +17.9$

TiB2 ♦ 10.0$ Si + 16.7$ Co by weight. This cermet body had the

-97- highest room temperature modulus of rupture of all compositions tested. It was also found to have good high temperature strength and good resistance to thermal shock.

Laboratory evaluation indicates that a suitable material for use in gas turbine blades has been developed. Blades should be prepared and tested under service conditions as a final test of the value of this cermet composition.

-98- CONCLUSIONS

Conclusions drawn from the results of this investigation ares

1. Titanium diboride and water react at 75°?. to form the min­

eral sassolite, E^O^‘SligO. Other products of this reaction are probably hydrogen and titanium dioxide.

2. During the processing of cermet materials which contained titanium carbide and titanium diboride milled in carbon tetrachloride, a reaction occurred in which titanium chloride compounds were formed.

A reaction between titanium carbide and carbon tetrachloride is thermodynamically feasible.

3. A composition consisting of 70% TiC ♦ 20% SiC + 10% B^C by weight was wetted only slightly under the conditions of these tests by cobalt, nickel, and NiAl at temperatures near the melting points of these metals. Chromium completely wetted this material at

3200°F. All of these metals wetted this mixed carbide at 3875°?*

4. Segregation of the constituents in sintered 70% TiC +

20% SiC + 10% B^C material was not caused by the reaction between titanium carbide and boron carbide as postulated by one investigator, but this segregation was the result of poor mixing technique.

5. The 70% TiC 20% SiC 10% B^C composition was not sintered dense alone or in combination with binder metals under the conditions of these tests.

6. The poor physical properties of sintered 70% TiC + 20% SiC+

10% B^G base composition alone and in combination with binder metals

-99- were caused by the presence of relatively high percentages of carbon and/or silicon carbide.

7. A body composed of 46.1$ TiC 23.9$ TiB2 + 13.3$ Si + 16.7$

Co sintered to a dense structure, and it had good oxidation resistance at 2000°F. during a three hundred hour test.

8. A cermet material (ill R B + Co) composed of 55.4$ TiC ♦

17.9$ TiB2 + 10»0$Si + 16.7$ Co also sintered dense and had good oxidation resistance during the three hundred hour test at 2000°F.

This body had the most favorable silicon to cobalt ratio of the compositions tested.

9. The above cermet composition had a room temperature modulus of rupture of 65,600 p.s.i. when properly sintered. It had moduli of rupture of 83,300 p.s.i. at 1600°F., 87,900 p.s.i. at 1800°F., and

65,800 p.s.i. at 2000°F.

10. The III.R B + Co material is resistant to thermal shock in small pieces. It also appears to possess high thermal conductivity.

11. Composition III S B ♦ Co is a promising material for use ' in gas turbine blades.

-100- BIBLIOGRAPHY

Books

1. Evans, Ulick R., Metallic Corrosion Passivity an d Protection, Edward Arnold and Co., London, 1938, p. 97.

2. Hansen, M., Per Aufbau der Zweistofflegierungen. Edwards Brothers, Inc., Ann Arbor, Mich., 1943, pp. 509-512.

3. Hodgman, Charles D., Editor, Handbook of Chemistry and Physics. Thirtieth Edition, Chemical Rubber Publishing Co., Cleveland, 0., 1946.

4. Mellor, J. W., A Comprehensive Treatise on Inorganic and Theoretical Chemistry, Longmans, Green and Co., London, 1925, Vol. 6, p. i W .

5. Staithells, C. J., Metals Reference Book, Interscience Publishers, Inc., New York, 1949, pp. 421-436.

Pamphlets and Periodicals

6. Accountius, Oliver E., "The High Temperature Oxidation Resistance of Systems Containing Titanium Carbide", Ph.D. Dissertation, The Ohio State University, Columbus, 0., 1951.

7. Baraduc-Muller, L., "Siliciures Metalliques-Action du Carbure de Silicium sur Quelques Oxydes Metalliques", Review of Metallurgy, Vol. 7, pp. 707-711, 1910.

8. Bennett, D. G., J. A. Nelso, T. A. Willmore, and R. C. IVomeldorph, "Refractory Bodies Composed of Boron and Titanium Carbides Bonded with Metals", Air Force Technical Report No. 6540, Wright-Patterson Air Force Base, Dayton, 0., 1951.

9. Boren, Bertil, "Rontgenuntersuchung der Legierungen von Silicium mit Chrom, Mangan, Kobalt und Nickel", Arkiv fur Kemi, Mineralogi, 0. Geologi, Bd. 11A, No. 10, S. 17,22,27, 1933.

10. Boren Bertil, Sven Stahl, and A. Westgren, "Kristallstruktur und Zuzammensetzung des rhombischen Kobaltsilicides", Z. Phycik, Chem. B., Bd. 29, S. 231-235, 1935.

11. Casagrande, Arthur, "The Hydrometer Method for Determination of Fineness of Soils", Julius Springer, Berlin, 1934.

12. Dawihl, W., and J. Hinnueber, "Structure of Cemented HardMetal Compositions", Kolloid Zeitschrift, Bd. 104, Nos. 2,3, S. 233-236, 1943, Translation No. 1545, Henry Brutcher, Altadena, Calif.

-101- 13. Deutsch, George C., Andrew J. Repko, and William G. Lidman, "Elevated Temperature Properties of Several Titanium Carbide Base Ceramals", National Advisory Committee for Aeronautics, Technical Note No. 1915, Washington, July, 1949.

14. Engle, E. W., "Cemented Carbides", Powder Metallurgy, (Edited by John Wulff), The American Society for Metals, Cleveland, 0., pp. 436-453, 1942.

15. Evans, Ulick R., "The Mechanism of the Formation of Films on Metals", (Pittsburgh International Conference on Surface Re­ actions), Corrosion Publishing Co., Pittsburgh, Pa., 1948.

16. Franssen, H., "Structure of Cemented Carbide Compositions", Archiv Fur Das Eisenhuttenwessen. Bd. 19, S. 79-84, 1948, Translation No. 2175, Henry Brutcher, Altadena, Calif.

17. Gangler, James J., Chester F. Robards, and James E. McNutt, "Physical Properties at Elevated Temperature of Seven Hot- Pressed Ceramics", National Advisory Committee for Aeronautics, Technical Note No. 1911, Washington, July, 1949*

18. Hamjian, H. J. and. W. G. Lidman, "Investigation of Bonding Between Metals and Ceramics - I - Nickel, Cobalt, Iron, or Chromium with Boron Carbide", National Advisory Committee for Aeronautics, Technical Note No. 1948, Washington, Sept., 1949*

19. Hoffman, Charles A., G. Mervin Ault, and James J. Gangler, "Initial Investigation of Carbide-Type Ceraraal of 80-Percent Titanium Carbide Plus 20-Percent Cobalt for Use as Gas-Turbine Blade Material", National Advisory Committee for Aeronautics, Technical Note No. 1836, Washington, March, 1949.

20. Honigschmid, M. Otto, "Sur le siliciure de zirconium ZrSi^ et silicure de titane TiSi^", C. R. Academy of Science. Paris. Vol. 143, PP. 224-226, 19067

21. Konrad, Howard E ., and Robert F. Stoops, "Preliminary Investi­ gation of Highly-Refractory Ceramic Bodies Suitable for Metal Impregnation", Master’s Thesis, The Ohio State University, Columbus, 0., 1951.

22. Lea, Arthur C., "VII - Silicon Carbide and Its Use as a Refractory Material", Transactions of the British Ceramic Society. Vol. 40, pp. 93-118, 1941.

23. Lebeau, M. Paul, "Sur un nouveau siliciure de cobalt", C. R. Academy of Science. Paris. Vol. 132, pp. 556-558, 1901.

24. Lebeau, M. Paul, "Sur les combinaisons du silicium avec le cobalt et sur un nouveau siliciure de ce metal", C. R. Academy of Science. Paris. Vol. 135, pp. 475-477, 1902.

-102- 25. Lefebvre, A., "Abrasives", Ind. Ceram., No. 382, pp. 13, 14, No. 383, p. 31, 1948.

26. Lenel, F. V., "Sintering in the Presence of a Liquid Phase", Metals Technology. Vol. 15, Technical Publication No. 2415, pp. 1-19, 1948.

27. Lewkonja, Kurt, "Uber die Legierungen des Kobalts mit Zinn, Antimon, Blei, Wisraut, Thallium, Zink, Cadmium, Chrom und Silicium", Z. Anorg. Ch., Bd. 59, S. 327, 1908.

28. Lidman W. G., and H. J. Hamjian, "Properties of a Boron Carbide - Iron Ceramal", National Advisory Committee for Aeronautics, Technical Note No. 2050, Washington, March, 1950.

29. Lustraan, Benjamin, "The Resistance of Metals to Oxidation at Ele­ vated Temperature", Metals Handbook. The American Society for Metals, Cleveland, 0., p. 223, 1949.

30. Meerson, G. A., G. L. Zverev and B. Ye. Qsinovskaya, "Investigation of Behavior of Titanium Carbide in Cemented Carbide Alloys", Zhurnal Prikladnoi Khimii, Vol. 13, No. 1, pp. 66-75, 1940, Translation No. 1642, Henry Brutcber, Altadena, Calif.

31. McBride, Clinton C., "NiAl - Titanium Carbide Cermets", Report No, 65, Project 341, The Ohio State University Research Foundation, Columbus, 0., Jan., 1951.

32. Norton, J. T., N. Blumenthal and S. J. Sindeband, "Structure of Diborides of Titanium, Zirconium, Columbium, Tantalum and Vanadium", Journal of Metals. Vol. 1, No. 10, pp. 749-751, Tran., 1949*

33. Nowotny, H., and G. Glenk, "Contributions to System Titanium Carbide - Tungsten Carbide", Metallforschung. Bd. 2, No. 9, S. 265-269, 1947, Translation No. 2211, Henry Brutcher, Altadena, Calif.

34* Oswald, Marcel, "Les alliages durs a base de carbures refractaires", Chimie et Industrie. Vol. 49, pp. 8-16, Jan. 1943*

35. Pilling, N. B., and R. E. Bedworth, "The Oxidation of Metals at High Temperatures," Journal of the Institute of Metals. Vol. 29, pp. 529-582, 1923.

3 6 . Redmond, John C., "Some Metallurgical Aspects of Cemented Carbides," Iron Age. Vol. 159, No. 5, pp. 42-45, 150, 1947.

37* Rhines, F. M., "Seminar on the Theory of Sintering," Metals Techno­ logy, Vol. 13, No. 5, pp. 1-18, Aug. 1946.

38. Rose, Kenneth, "Cemented Carbides," Materials and Methods. Vol. 29, No. 2, pp. 73-84, Feb. 1949.

-103- 39. Sabol, Frank P., "Sintered A 1 ^ 3 and Combinations of Ald03 and SiC," Report No. 33, Project 252, The Ohio State University Research Foundation, Columbus, 0., Oct. 1947.

40. Sindebrand, S. J., "Properties of Chromium Boride and Sintered Chromium Boride", Journal of Metals. Vol. 1, No. 2, pp. 198-202, Feb. 1949.

41. Skaupy, F., "Cemented Carbides Considered from the Viewpoint of Dispersoid Chemistry - Communication 1," Kolloid Zeitschrift, Bd. 98, S. 92-95, 1942, Translation No. 1413, Henry Brutcher, Altadena, Calif.

42. Skaupy, F., "Cemented Carbides Considered from the Viewpoint of Dispersoid Chemistry - Communication 2", Kolloid Zeitschrift. Bd. 102, S. 269-271, 1943, Translation No. 1558, Henxy Brutcher, Altadena, Calif.

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4 6 . Whitman, M. J., and A. J. Repko, "Oxidation of Titanium Carbide Base Ceramals Containing Molybdenum, Tungsten, and Cobalt," National Advisory Comnittee for Aeronautics, Technical Report No. 1914, Washington, July, 1949*

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Patents

49. Heyroth, Albert H., "Impregnated Silicon Carbide Article and the Manufacture Thereof", U. S. Patent No. 1,906,963, May 2, 1933, (to Globar Corporation, Niagra Falls, N. I.).

-10 V AUTOBIOGRAPHY

I, Robert Franklin Stoops, was born in Winona, West Virginia,

June 16, 1921. I received my secondary school education in the

public schools of the city of Staunton, Virginia. My undergraduate

training was obtained at North Carolina State College of Engineering

and Agriculture, from which I received the degree Bachelor of Ceramic

Engineering in 1949* From The Ohio State University, I received the

degree Master of Science in 1950. While completing the requirements

for this degree and for the degree Doctor of Philosophy, I was em­

ployed as a Research Fellow by The Ohio State University Research

Foundation.

-105-