Process for producing dispersed particulate composite materials

This invention is directed to a process for forming noninterwoven dispersed particulate composite products. In one case a composite multi-layer film product comprises a substantially noninterwoven multi-layer film having a plurality of discrete layers. This noninterwoven film comprises at least one discrete layer of a first material and at least one discrete layer of a second material. In another case the first and second materials are blended together with each other. In either case, the first material comprises a metalloid and the second material a metal compound. At least one component of a first material in one discrete layer undergoes a solid state displacement reaction with at least one component of a second material thereby producing the requisite noninterwoven composite film product. Preferably, the first material comprises silicon, the second material comprises Mo.sub.2 C, the third material comprises SiC and the fourth material comprises MoSi.sub.2.

In U.S. Pat. No. 4,454,015, an anode electrode composition having an 
improved electrical conductivity is provided by contacting a combination 
of metallic and metallic-oxides, oxygen-containing compounds, and an 
elevated temperature resulting in a displacement reaction to form an 
interwoven network of metal oxides and metal alloy. Metal oxides include 
oxyborides, oxynitrides, and oxyhalides. Metal borides, nitrides, 
carbides, halide and sulfides are also said to be within the term metal 
compounds. In U.S. Pat. No. 4,465,577, cermet films are formed which 
include a substrate and multiple, ultra-thin alternating layers of 
conductive and non-conductive material on the substrate. The layers are 
created using RF sputtering techniques to form alternating layers of 
discontinuous island of particles from each of the different materials. 
Each of layers is on order of 20-50 angstroms thick. The substrate is 
silicon, silicon dioxide or aluminum trioxide. In examples of the cermet 
layers are alternating layers of silicon dioxide in either gold or 
chromium. 
Solid state displacement reactions are diffusional phase transformations 
whereby two or more elements or compounds react to form new product 
compounds that are more thermodynamically stable than the starting 
reactants. Prior art interwoven microstructures can be achieved using 
displacement reactions. This particular work grew out of studies on 
displacement reactions involving Ni/CuO.sub.2, Ni/NiO, Co/Cu.sub.2 O, 
Fe/Cu.sub.2 O, Fe/NiO (See References 1-4 below). Others, such as provided 
in References 5-14 below, have used a related approach termed reaction 
sintering as a route to ceramic composites. However, this has met with 
limited success in producing high quality structural materials due to the 
lack of predictability and understanding of reaction mechanisms and 
reaction morphologies (see References 3, 15, 16 below). It is difficult to 
understand the formation kinetics for these solid state displacement 
reactions. For any specific set of potential reactants one cannot predict 
with certainty the ability and extent of the displacement reaction due to 
a lack of critical materials data required for making that determination. 
References 
1. R. A. Rapp, A. Ezis and G. J. Yurek, Met. Irans 4, 1283 (1973). 
2. G. J. Yurek, R.a. Rapp and J. P. Hirth, Met. Trans. 4, 1293 (1973). 
3. G. J. Yurek, R. A. Rapp and J. P. Hirth, Met. Irans. 10A, 1473 (1979). 
4. C. Tangchitvittaya, J. P. Hirth and R. A. Rapp, Met. Trans. 13A, 585 
(1982). 
5. J. S. Haggerty and U. M. Chiang, Ceram. Eng. Sci. Proc. 11 (7 8)757 (19 
90). 
6. N. Claussen and J. Jahn, J. Amer. Ceram, Soc. 63, 228 (1980). 
7. M. R. Ansedau, C. Leblud, F. Cambier, J. Mater. Sci. Lett. 2, 366 
(1983). 
8. S. Yangyun and R. J. Brook, Ceram, Int. 9, 39 (1985) 
9. P. Boch and J. P. Giry, Mater, Sci., and Eng. 71, 39 (1985) 
10. A. Mocellin and G. Bayer, J. Mater, Sci. 20 3691 (1985) 
11. P. Pena, P. Miranza, J. S. Moya and S. DeAza, J. Mater. Sci. 20, 2011 
(1985). 
12. P. Miranza, P. Pena, J. S. Moya and S. De Aza, J. Mater. Sci, 20, 2702 
(1985). 
13. G. Bayer and A. Mocellin, Rev. Chimie Miner, 23, 80 (1986). 
14. P. Miranza, P. Pena, S. DeA.a, J. S. Moya, J. MaRincon, and G. Thomas, 
J. Mater Sci, 22, 2987 (1985). 
15. M. Martin and H. Schmalzried, Ber. Bunsenges. Phys. Chem, 89, 124 
(1985). 
16. M. Ackhaus-Ricoult and H. Schmaizried, Bayer. Bunsenges, Phys. Chem. 89 
1323 (1985). 
17. R. D. Doherty, Physical Metallurgy, 3rd edition. R. W. Cahn and P. 
Haasen, eds., p. 933, North-Holland, New York (1983). 
18. E. W. Lee, J. Cook. A. Khan, R. Mahapatra, and J. Waldman, J. of Metals 
3, 54 (1991). 
19. E. J. J. Vanloo, F. M. Smet, G.d. Rieck, and G. Verspui, High 
Temp.-High Press. 14, 25 (1982). 
20. D. II. Carter. W. S. Gibbs, and J. J. Petrovic, Ceramic Materials and 
Components for Engines, V. J. Tennery, ed., p. 977, the American Ceramic 
Society, Westerville, Ohio (1989). 
SUMMARY OF THE INVENTION 
This invention relates to a process for forming a substantially 
noninterwoven dispersed particulate composite material. 
In the case of a dispersed particulate composite multi-layer product, a 
substantially noninterwoven multi-layer structure is first formed having a 
plurality of discrete layers. This noninterwoven film comprises at least 
one discrete layer of a first material and at least one discrete layer of 
a second material. Typically, the first material comprises a metalloid 
compound, preferably silicon, and second material comprises a metal 
compound, such as Mo.sub.2 C. 
These first and second materials preferably have an average particle size 
of not more than about one micron, more preferably not more than about 
0.75 microns, and most preferably not more than about 0.50 microns. This 
small particle size and the close proximity of the first and second layers 
minimizes the diffusion length and controls the composite phase separation 
distance so that a solid state displacement reaction can effectively and 
efficiently occur. The typical thickness of the layers of the first and 
second material is not more than about 5.0 microns, preferably not more 
than about 3.0 microns, and more preferably not more than about 1.0 
microns. 
The discrete layers of the dispersed particulate composite, which are 
substantially unreactive with each other at ambient temperature, are 
preferably arranged in alternating adjacent layers. The materials in these 
alternating adjacent layers undergo a solid state displacement reaction to 
form a multi-component third material and a multi-component fourth 
material in which particles of the third material are dispersed within a 
matrix of the fourth material, thereby producing said composite product. 
In order to effect the solid state displacement reaction of the present 
invention, at least one of the first material or second material should 
have a multi-component structure. A multi-component structure means that 
the material includes at least two chemical elements. Examples of such 
multi-component structures are Mo.sub.2 C and B.sub.4 C. 
A specific reaction involving the subject process can relate to a first 
material comprising silicon, a second material comprising Mo.sub.2 C, 
forming third and fourth materials comprising SiC and MoSi.sub.2. The 
solid state displacement reaction can comprise the following chemical 
equation: 
EQU Mo.sub.2 C+5Si.fwdarw.2MoSi.sub.2 +SiC 
Another reaction involving the subject process relates to first material 
comprising titanium, a second material comprising B.sub.4 C, and the 
production of a product comprising TiB.sub.2 and TiC. This solid state 
displacement reaction comprises the following chemical equation: 
EQU 3Ti+B.sub.4 C.fwdarw.2TiB.sub.2 +TiC 
Preferably, as evidenced by the above solid state displacement reactions, 
the process of this invention can comprise reactants wherein neither of 
the first and second materials are oxygen-containing. 
This process also contemplates solid displacement reactions in which the 
first material comprises Ni.sub.2 Al.sub.3, the second material comprises 
NiO, and the third and fourth materials product comprises Al.sub.2 O.sub.3 
and NiAl.sub.2 O.sub.4, as well as reactions in which the first material 
comprises sappharine (4MgO.5Al.sub.2 O.sub.3.2SiO.sub.2), the second 
material comprises alumina (Al.sub.2 O.sub.3), and the thin 
ceramic-containing composite film product comprises mullite 
(2SiO.sub.2.3AlO.sub.3) and spinel (MgAl.sub.2 O.sub.4). 
The noninterwoven dispersed particulate composite material is then heated 
to a temperature wherein at least one component of a first material 
undergoes a solid state displacement reaction with at least one component 
of a second material. In this way, the composite product is formed. 
In the process of present invention it is preferred that the particles of 
the third material comprise platelets and more specifically that the 
particles of the third material are uniformly distributed throughout the 
fourth material. Moreover, it is also preferred that the first material 
and the second material interdiffuse, and the alignment of said particles 
of the third material with respect to the fourth material is in the 
direction of said interdiffusion of the first and second material. 
Typically, the first material and the second material comprise a diffusion 
couple. Finally, the first material and the second material can comprise 
powders which are formed into said discrete layers. 
In another form of this invention, the process for forming a substantially 
noninterwoven composite product comprises blending together a first powder 
material comprising a metalloid compound and a second powder material 
comprising a metal compound to form a blended powder as described above. 
The blended powder is heated wherein at least one component of a first 
material and at least one component of a second material undergoes a solid 
state displacement reaction to form particles of a third material and a 
fourth material. As before, these particles of the third material are 
dispersed within a matrix of the fourth material to produce the 
substantially noninterwoven dispersed particulate composite product. 
The foregoing and other objects, features and advantages of the invention 
will become more readily apparent from the following detailed description 
of a preferred embodiment of the invention which proceeds with reference 
to the accompanying drawings.

DETAILED DESCRIPTION OF PREFERRED EMBODIMENTS 
Solid state displacement reactions can produce in situ intermetallic matrix 
composites in a process where an intermetallic phase(s) and a potential 
reinforcing phase(s) are grown together during a solid state reaction. 
Dispersed microstructures, important for desirable dispersed particulate 
composite properties, have been produced by means of displacement reaction 
processing techniques. Such composites are MoSi.sub.2 reinforced with SiC 
particles. Strength in bending and chevron-notch fracture toughness have 
been determined as a function of temperature and measured properties 
compare favorably with composites produced by other means. 
EXAMPLE 1 
Dispersed particulate composite materials were formed employing the process 
of the present invention. 
Vacuum hot-pressed Mo.sub.2 C (99.5% purity) and pieces cut from Si single 
crystal wafers (99.99% purity) were used for the diffusion couples. 
Diffusion couples were prepared using cut and polished pieces roughly 
0.5.times.0.5 cm.sup.2 and 0.3 cm thick. Several couples were made using 
Si pieces only 0.025 cm thick. These were placed in a Mo holder with 
pieces of Al.sub.2 O.sub.3 and Mo foil. The thermal expansion of this 
material sandwich in the holder was used to hold the Mo.sub.2 C and Si 
pieces in contact at elevated temperatures. The couples were heated in a 
vacuum furnace operated at pressures less than 10.sup.-6 torr. 
Powders of Mo.sub.2 C (-335 mesh, 99+% purity) and Si (-335 mesh, 99.99% 
purity) were used for the vacuum hot-pressed materials. The powders were 
blended in a 5:1 Si:Mo.sub.2 C ratio using a vibratory ball mixer and 
hot-pressed at 27.5 MPa using graphite dies under a vacuum of about 
10.sup.-4 torr. The hot-pressing temperature cycle was 1350.degree. C. for 
2 hours followed by 1 hour at 1700.degree. C. A hot-press die diameter of 
2.2 cm was used with 10 to 15 grams of blended powders for the specimens 
produced here. 
Diffusion couples were reacted at 1200.degree. C. for 16 hours and at 
1350.degree. C. for 20 hours. The reaction at 1200.degree. C. produced an 
interesting microstructure of SiC ribbons in a MoSi.sub.2 matrix (FIG. 1). 
The following displacement reaction occurred: 
EQU Mo.sub.2 C+5Si.fwdarw.2MoSi.sub.2 +SiC 
The areal fraction of the SiC particles was determined to be 
30%. Although the SiC lamellae exhibit a rod-like appearance in this view, 
the basic shape is more ribbon-like (FIG. 1a). FIG. 1b shows the 
microstructure from a section in which the original interface is parallel 
to the paper. The narrow dimension of the SiC particles was less than 1 
micron and they had a length of 10 microns. Some of the lamellae appear to 
be blocky particles but this could be a result of cutting the ribbon-like 
lamellae at an angle and thereby increasing the effective dimensions. The 
diffusion couple reacted at 1350.degree. C. revealed similar 
microstructures except that indications of spheroidizing of the SiC 
particles were observed. 
Scanning electron microscopy (SEM) was performed for phase and morphology 
identification. Energy dispersive X-ray (EDX) analysis showed no 
indication of Mo in the SiC, at least to the accuracy of the analysis. 
There was also no indication of carbon in the MoSi.sub.2 ; however, the 
EDX analysis for carbon is not very sensitive. In SiC, which is 50% 
carbon, the carbon peak is very small. 
Microhardness indents were used to determine the hardness and crack 
propagation behavior of this material. Cracking was observed around the 
hardness indent but was not typical of classically brittle behavior. There 
was rubblelizing at the indent rather than single cracks emanating from 
the indent corners as would be expected for a brittle ceramic. There were 
indications of crack deflection at the SiC lamellae and crack 
wake-bridging (FIG. 2) that are quite similar in appearance to whisker 
reinforced ceramic materials. The microhardness was 1292 HV.sub.1000 (12.8 
GPa) in the reacted region. 
Composites fabricated from the blended powders and hot-pressing using the 
schedule of 2 hours at 1350.degree. C. and 1 hour at 1700.degree. C. were 
similar to those produced in the diffusion couples with the exception of 
the morphology of the SiC phase (FIG. 3a) and of the observation of 
Mo.sub.5 Si.sub.3 regions (FIG. 3b), which were identified using EDX. The 
SiC particles were more equiaxed, most likely due to spheroidization of 
the platelets, but were uniformly distributed throughout the material. 
Larger regions, up to 20 .mu.m in diameter, of Mo.sub.5 Si.sub.3 are 
dispersed throughout the material. Xray diffraction (XRD) reveals strong 
MoSi.sub.2 peaks, SiC as a second phase, and some very faint peaks that 
could not be indexed but likely belong to the Mo.sub.5 Si.sub.3 identified 
in the SLM analysis. The SiC particles are smaller than 1 .mu.m in 
diameter. 
The composite appeared to be near full density from observing polished 
surfaces in the SEM. The density was measured to be 5.53 g/cm.sup.1. An 
estimated theoretical density for a MoSi.sub.2 /SiC (30 vol % SiC) 
composite is 5.35 g/cm.sup.3. A hardness of 1440.+-.12 HV.sub.1000 
(14.2.+-.0.1 GPa) was measured and an indentation fracture toughness of 
8.7.+-.0.1 MPa.sqroot.m was calculated for this composite using the 
observed median cracking under 15 kg Knoop indents. The cracking pattern 
was irregular enough (rubbelizing behavior) to suggest a large uncertainty 
in this toughness value. Crack interactions with the SiC particles, such 
as crack deflection and crack-wake bridging, were observed similar to that 
shown in FIG. 2. 
The formation of SiC particles having a plate-like morphology in MoSi.sub.2 
from the displacement reaction between Mo.sub.2 C and Si is quite 
extraordinary, particularly since the SiC volume fraction is nearly ideal 
for a whisker-reinforced composite. While the volume fraction can be 
estimated from the balanced reaction, the SiC morphology is difficult to 
predict with the available models and data. We expect the platelet size to 
depend on the reaction temperature, as for other discontinuous 
precipitation reactions (see Reference 17 above). In general, this control 
over the particle size is one of the strong points of displacement 
reaction synthesis of in situ composites. Comparison of the microstructure 
obtained here (FIG. 3) with that of the Martin Marietta XD.TM. process 
(see Reference 18 above) clearly shows that displacement reaction 
processing provides a more uniform dispersion of SiC in MoSi.sub.2 and a 
more uniform size dispersion as well. The ability to control the size and 
morphology of the SiC particles, while not completely demonstrated here, 
gives displacement reaction processing important advantages over some of 
these other synthesis routes. 
The alignment of the SiC platelets in the direction of diffusing species in 
the diffusion couple suggests a cooperative precipitation reaction and it 
is likely that SiC is being formed by the cooperative rejection of some 
species, probably carbon, during interdiffusion of Si and Mo.sub.2 C to 
form MoSi.sub.2. This view receives further support on examining the 
Mo--Si--C ternary phase diagram (see Reference 19 above), 1200.degree. C. 
cross-section (FIG. 4). That work reports the formation of a ternary 
Mo.sub.5 Si.sub.3 C phase formed from diffusion couples between Mo and 
SiC. A reaction pathway between Mo.sub.2 C and Si would probably include 
the Mo.sub.2 Si.sub.3 C phase, which could provide a source of carbon from 
the reaction Mo.sub.5 Si.sub.3 C.fwdarw.Mo.sub.5 Si.sub.3 +C. Further 
interdiffusion of Si is suggested to result in the formation of SiC and 
MoSi.sub.2. 
The powder reaction during hot-pressing is expected to follow the same 
pathway, resulting in SiC platelets randomly oriented in the hot-pressed 
body following the reaction at 1350.degree. C. The distinction is the 
further coarsening of the SiC after 1 hour at 1700.degree. C. The SiC 
volume fraction and interparticle spacing appears to be similar for the 
diffusion couple and the hot pressed body. Observations of hot pressed 
bodies at 1350.degree. C. without the 1700.degree. C. densification step 
revealed that the reaction was completed in two hours at 1350.degree. C., 
which is consistent with the short diffusion lengths in the blended 
powders. The presence of Mo.sub.5 Si.sub.3 in the hot pressed material 
suggests that insufficient Si was used or that inadequate mixing occurred 
during powder blending. 
The interaction of Knoop indentation cracks with the SiC platelets 
indicates that some measure of improved fracture toughness can be expected 
from these materials. Crack deflection along the MoSi.sub.2 /SiC 
interfaces means that this interface is weaker than the SiC particles 
which is required for crack-wake bridging and crack deflection processes 
to occur. The SEM observations do not show any indication of glassy phases 
at these boundaries (see Reference 20 above). The indentation fracture 
toughness of 8.7 MPa.sqroot.m is consistent with that measured for SiC 
whisker-reinforced MoSi.sub.2 (see Reference 20 above) but the indentation 
data obtained here, while indicative of increased toughness compared to 
monolithic MoSi.sub.2, is clouded by the non-ideal Knoop cracking patterns 
observed. 
Chevron-notched bend bars are being prepared from larger hot-pressings to 
obtain better data. Bend strength as a function of temperature will also 
be explored later using four point bend bars. It is anticipated that this 
composite will have excellent oxidation resistance based on the behavior 
of other MoSi.sub.2 /Sic composites (see Reference 18 above) and good 
creep strength based on the uniform dispersion of SiC particles. 
The solid state displacement reaction between Mo.sub.2 C and Si was used to 
synthesize a MoSi.sub.2 /SiC composite in situ. A diffusion couple 
processed at 1200.degree. C. proceed SiC platelets in a MoSi.sub.2 matrix. 
The SiC platelets had an areal fraction of about 30% and were nominally 1 
.mu.m wide and 10 .mu.m long and were aligned in the direction the 
diffusion species. Cracks induced from Knoop indentations exhibited 
deflection along the MoSi.sub.2 /SiC interfaces and crack-wake binding was 
observed. Composites made by blending Si and Mo.sub.2 C powders and 
hot-pressing at 1350.degree. C. for 2 hours followed by 1 hour at 
1700.degree. C. consisted of about 30 vol % SiC particles uniformly 
dispersed in a MoSi.sub.2 matrix. The particles were about 1 .mu.m in 
diameter and appeared to be spheroidized versions of the SiC platelets 
obtained at 1200.degree. from the diffusion couples. The hot pressed 
material had a density of 5.53 g/cm.sup.3 and an indentation fracture 
toughness of 8.7 MPa.sqroot.m. 
Composites were made by VHP of blended Si and Mo.sub.2 C powders at 
1350.degree. C. for 2 hrs. followed by densification at temperatures in 
the range of 1600.degree. C. to 1800.degree. C. undergo the following 
displacement reaction as determined by x-ray diffraction (XRD) and 
quantitative metallography: 
##STR1## 
This is the basic, stoichiometric reaction which, based on the presence of 
only a trace amount of the ternary Mo.sub.5 Si.sub.3 C phase, proceeds 
nearly to completion at 1600.degree. C. The .DELTA.G.sup.298.sub.Total for 
the above reaction is -310 kJ/mol and approximately 25% volume shrinkage 
occurs. Consideration of the ternary Mo--Si--C phase diagram and further 
analysis of diffusion couples between Mo.sub.2 C and Si indicate the 
following reaction sequence and characteristics: 1) Si is the fastest 
diffusing species and the ternary Mo.sub.5 Si.sub.3 C phase forms 
initially in the Mo.sub.2 C phase, followed by the MoSi.sub.2 phase, 2) 
the SiC is observed to grow at the interface between the ternary Mo.sub.5 
Si.sub.3 C phase and MoSi.sub.2, but within the ternary phase, 3) the SiC 
growth direction is into the ternary phase, 4) the initial SiC morphology 
is plate-like with an aspect ratio of .about.20, 5) the SiC plates undergo 
a Rayleigh instability and pinch-off into discrete particles approximately 
1 .mu. m in diameter at longer times, 6) the grain size of the MoSi.sub.2 
phase is less than 1 .mu.m, and 7) the composites, after final 
consolidation, occupy compositions within a three-phase triangle given by 
MoSi.sub.2 --SiC--Mo.sub.5 Si.sub.3 C. 
Two additional reactions were investigated by adding C powders to explore 
other compositions within the three-phase field: 
VHP-3: Mo.sub.2 C+5Si+0.3C.fwdarw.1.8MoSi.sub.2 +1.25SIC+0.05 Mo.sub.5 
Si.sub.3 C 
VHP-4: Mo.sub.2 C+5Si+0.6C 1.6MoSi.sub.2 +1.50SIC+0.1 Mo.sub.5 Si.sub.3 C 
Adding C reduces the MoSi.sub.2 grain size, increases the porosity for 
equal VHP times, and moves the final composition into the interior of the 
three-phase triangle towards C as expected. It is clear that other 
compositions within the triangle can be achieved by control of the 
relative amounts of Mo, Si, and C. 
Processing by means of VHP of blended powders results in a 
dispersed-particulate phase composite due to the pinch-off instability of 
the SiC plates at 1600.degree. C. to 1800.degree. C. Representative 
polished cross-sections imaged by scanning electron microscopy (SEM) 
reveal the differences between the compositions and show the uniform 
dispersion of the three phases in the composites. Photomicrographs of 
fracture surfaces reveal the MoSi.sub.2 grain size and show that the SiC 
particles lie mainly on the grain boundaries and at triple points. This 
suggests that the MoSi.sub.2 grain size is controlled by the SiC volume 
fraction and spacing. 
Mechanical properties (strength and toughness) of the composite materials 
as a function of temperature and composition suggest that both strength 
and toughness are highest for the material made by the above reaction 
consolidated at 1800.degree. C. This material has the largest grain size, 
smallest volume fraction of SiC and Mo.sub.5 Si.sub.3 C, and least 
porosity. Adding C reduces both fracture strength and toughness, and does 
not increase elevated temperature strengths. An apparent ductile-brittle 
transition (DBT) occurs at about 1000.degree. C. for all these materials. 
Above this DBT, with the exception of the material with the highest 
C-content and smallest grain size, the materials behave identically with 
respect to strength decrease as a function of test temperature, suggesting 
that strength above the DBT is controlled by deformation within the 
MoSi.sub.2, which is the continuous phase. The fact that the DBT is 
.about.1000.degree. C. rather than 1300.degree. C. is likely due to the 
fine grain size and increased contribution of grain boundary sliding to 
the deformation of the material. 
Bend strength and chevron-notched fracture toughness were determined for 
all materials as a function of temperature in a self-aligning, SiC 4-point 
bend fixture having a 40 mm lower span and a 20 mm upper span. The 
specimen sizes were all 4 mm.times.4 mm.times.50 mm (nominal). The chevron 
notches were cut with a 60.degree. included angle such that the tip of the 
chevron was approximately flush with the specimen surface. All chevron 
notches were measured after fracture testing, however, and a geometry 
factor was calculated for each specimen. In addition, chevron-notched 
specimens of AD995 alumina and Pyrex glass were tested at room temperature 
to calibrate the calculated geometry factor. Seven Pyrex specimens and 6 
alumina specimens were tested, and the average fracture toughness values 
obtained were 0.81.+-.0.11 MPa.sqroot.m and 3.73 .+-.0.12 MPa.sqroot.m, 
respectively, using the appropriate calculated geometry factor Y*.sub.min. 
These values compare favorably with those reported in the literature. 
All bend tests were conducted in air at a crosshead speed of 
8.5.times.10.sup.-4 mm/s in a MoSi.sub.2 -element vertical tube furnace. A 
SiC support tube held the bend fixture at the furnace mid-plane. Specimen 
deflections were measured at the mid-point of the bend bar by means of an 
alumina rod attached to a strain gage extensometer. All bend data were 
corrected for the fixture compliance. 
The elevated temperature strength of the material with the highest C 
content suggests that deformation for this material above the DBT is 
controlled by a different process since the rate of strength decrease with 
temperature is lower. It may be that the MoSi.sub.2 is no longer 
continuous given the higher volume fractions of SiC and Mo.sub.5 Si.sub.3 
C. This should result in a material more resistant to deformation at 
elevated temperatures since SiC does not deform at 1000.degree. C., but 
nothing is known of the mechanical properties of the Mo.sub.5 Si.sub.3 C 
phase. 
The composite material's increased toughness, in comparison to pure 
MoSi.sub.2 which has a toughness of about 3 to 5 MPa.infin.m, appears to 
originate with SiC-particle bridging in the crack wake and crack 
deflection at the SiC/MoSi.sub.2 interfaces. The failure mode of the 
MoSi.sub.2 phase appears to be transgranular cleavage at room temperature, 
and many examples of SiC particle pullout can be seen in the fractographs. 
The highest toughness would, therefore, be expected for the material with 
the largest volume fraction of SiC within the MoSi.sub.2 grains, which 
would be the material with the largest grain size. It is encouraging that 
the fracture toughness is nominally equivalent to whisker- and 
particle-reinforced MoSi.sub.2 materials, which suggests that in situ 
techniques can be used in place of expensive and hazardous whiskers to 
obtain significant property improvements. 
A dispersion of SiC within MoSi.sub.2 is produced wherein the SiC phase is 
not interconnected. Interdiffusion of Si into Mo.sub.2 C acts to 
destabilize planar growth interfaces to produce a dispersed structure. In 
the MoSi.sub.2 /SiC system, the true displacement reaction occurs between 
the ternary Mo.sub.5 Si.sub.3 C phase and the Si phase to form MoSi.sub.2 
+SiC. The reaction between Si and Mo.sub.2 C to give Mo.sub.5 Si.sub.3 C 
will tend to produce a stable layer of Mo.sub.5 Si.sub.3 C, which is 
observed, because of the equilibrium join between the Mo.sub.2 C and 
Mo.sub.5 Si.sub.3 C. A layer of MoSi.sub.2 is then observed to form 
between the Si and the Mo.sub.5 Si.sub.3 C phase, and the SiC then forms 
at the MoSi.sub.2 /Mo.sub.5 Si.sub.3 C interface within the Mo.sub.5 
Si.sub.3 C phase. The SiC grows initially as platelets into the Mo.sub.5 
Si.sub.3 C phase and is later entrained within the MoSi.sub.2 phase as the 
Mo.sub.5 Si.sub.3 C decomposes into MoSi.sub.2 +SiC. 
EXAMPLE 2 
The solid state displacement reaction of Example 1 was repeated by heating 
the diffusion couple formed by equimolar amounts of pure Ni.sub.2 Al.sub.3 
and Ni.sub.4 O, in a molybdenum holder with Al.sub.2 O.sub.3 spacers, at a 
vacuum of &lt;l.times.10.sup.-6 torr and at a 1200.degree. C. for 16 hours. 
Expansion of the Al.sub.2 O.sub.3 relative to the molybdenum provided a 
constant compression of the couple during the heating process. After 
heating, the couple was sectioned and the microstructure of the reaction 
zone was analyzed. 
The reaction was determined to produce NiAl, Al.sub.2 O.sub.3 and 
NiAl.sub.2 O.sub.4. The final product was analyzed as described in Example 
1. Analysis of the reaction zone between the NiO and the Ni--Al alloy 
revealed a pure nickel layer, an Al.sub.2 O.sub.3 layer, and very large 
zone of NiAl with a dispersion of Al.sub.2 O.sub.3 particles. Volume 
fraction of the oxide particles was approximately 25% in the middle of the 
reaction zone, but was less near the edge of the zone. It was determined 
that Al.sub.2 O.sub.3 was dispersed in the NiAl phase. 
EXAMPLE 3 
Composites have also been synthesized by displacement reactions between 
NiAl and NiO powders to produce a composite material consisting of NiAl or 
NiAl+Ni.sub.3 Al+g-Ni and Al.sub.2 O.sub.3. Powders of NiAl (d&lt;45 .mu.m, 
99.5% purity) and NiO (d&lt;45 .mu.m, 99% purity) were blended for the VHP 
powder compacts. The powders were blended in 4:1 and 3:2 NiAl:NiO mol 
ratios in a vibratory ball mixer and hot-pressed at 27.5 MPa in graphite 
dies under a vacuum of about 10.sup.-2 Pa. A hot-press die diameter of 
7.62 cm was used with .about.185 g of blended powders. A hot-pressing 
temperature of 1300.degree. C. for 3 h was used for the NiAl-based VHP 
compacts. 
In the Ni--Al/NiO system, the planar or layered interfaces are apparently 
stabilized during interdiffusion and the interconnected network of 
Al.sub.2 O.sub.3 cells develop around the Ni--Al intermetallic regions. 
This is in accord with reaction products observed after diffusion took 
place at 1200.degree. C. in a couple between NiO and Ni.sub.2 Al.sub.3. 
The products consisted of discrete and stable layers of Ni/NiAl.sub.2 
O.sub.4 /Al.sub.2 O.sub.3 /NiAl. It would be expected that any .gamma.-Ni 
that formed during the reaction would be located at the Al.sub.2 O.sub.3 
/NiAl interface. 
Having described and illustrated the principles of the invention in a 
preferred embodiment thereof, it should be apparent that the invention can 
be modified in arrangement and detail without departing from such 
principles. I claim all modifications and variation coming within the 
spirit and scope of the following claims.