IL44318A - Alumina carbide ceramic material - Google Patents

Abstract

The material made from alumina and titanium carbide has a mean transverse breaking strength of at least 67.9 + 14.7 kg/mm<2>, a mean Rockwell A hardness of 93.8 and a mean Knoop hardness of 3477. It is produced by sintering a powder mixture of 70 % of alumina and 30 % of titanium carbide while the pressure is varied during the heating under specified conditions. The material can be used as refractory material and for the production of cutting tools for metals. [GB1465152A]

Description

44318/2 "V Alumina carbide ceramic material THE BABCOCK & WILCOX COMPANY c. 42323 This invention relates to a method of producing a refractory material and material produced by such a method. umina (AI2O3) and alumina compounds have been used for high temperature and high strength purposes for many years. For example, in refractory applications and in metal working tools that are subjected to high speeds and great wear, these materials have found widespread industrial acceptance.

It appears, moreover, that the strength of this material is in some manner related to its density and crystal size, the more dense and small crystal structures providing stronger and more durable tools. Consequently, there is a great deal of emphasis "on producing ceramic cutting tools with these characteristics. When used as a cutting edge, however, alumina occasionaly fractures. In general, these fractures seem to be related to the presence of relatively large alumina crystals, or "grains", in an essentially small crystal or "fine" grain structure. Thus, much of the alumina research effort has been directed to the more specific development of techniques for large-scale production of a high density material with a uniformly fine grain structure.

The crystal growth that .occurs when the raw powder material is heated to coalesce or is "sintered") often is retarded through the addition of magnesium oxide (MgO) in an amount of 0.5% or less. This heating can be accomplished in a vacuum furnace that raises the material temperature to a 1400° to 1550° range. Processes of this sort have been reported to provide a material that has a crystal size of approximately 2 to 3 microns. To attain this result, however, heating times in excess of four hours during sintering are required.

In the interest of efficiency and production economy, it is clear that a reduction in heating time is desirable, especially a more uniformly fine grain structure. An example of such efforts, U.S. Patent Specification No. 3,580,708 outlines a process employing fixed temperature and pressure parameters throughout the entire sintering period. Because of the tendency for alumina tools to fracture, there also is a need for a technique to produce the even smaller crystal sizes that lead to greater strength.

According to the present invention, there is provided a method for producing a refractory material comprising the steps of ball milling a refractory carbide, nitride or oxide powder, mixing' sinterable ametal oxide powder and said carbide, nitride or oxide powder, compressing said mixed powders to eliminate entrapped gas , heating said compressed powders at a first rate of 400° to 1000°C per minute, applying a constant physical pressure to said powders while the powders are being heated at said rate, applying an increased physical pressure to said powders upon reaching an onset of powder shrinkage to attain a maximum hot process press re, heating said powders at a second rate lower than said first rate to attain a maximum temperature while applying said increased pressure, maintaining-said maximum hot process pressure and said maximum temperature for two to six minutes, and releasing the pressure and cooling the material at the end of said period and prior to bloating.

The invention will now be further described by way of example with reference to the accompanying drawings, in which: Figure 1 is a schematic graph of ram displacement versus time to illustrate a "break away point"; and Figure 2 is an array of graphs that show pressure, temperature, density and breakaway point as function of time for a number of materials.

Figure 1 graphically illustrates features of examples of the invention expressed in terms of the movement or displacement of to remove any entrapped gases in the powder between time 0 and t After time t^ and before time t2, the application of heat to the pre- " compressed powder leads to a thermal expansion displacement 12 of the ram. This step in the process is terminated by a "break away point" 13 at the time t2« This "break away point" is characterized by a change from the expansion of the precompressed powder to a contraction 14 that commences as sintering begins. The contraction culminates at time tj. The time t3 is a time of maximum densification and coalescense of the sintered powder. The further application of heat after time t produces excessive grain growth or "bloating" 16 as indicated by the increase in ram displacement. It is at this time t3, before the material starts bloating, that the process is terminated.

Figure 2 is a graphic representation of pressure, sintered product temperature,density, and "break away points" as a function of time for the following materials: Billet Diameter Material 1/4" uo2 1" A1203 5" A1203 1" A1203 5" A1203 For purposes of orientation between Figures 1 and 2 the initial time, zero, of Figure 2 corresponds to the time t^ in Figure 1.

The pressure "history" 20 for all of these materials is bounded by straight line segments that identify a pressure increase, a step function from the initial pressure to the maximum hot process pressure that is maintained throughout the .remainder of the process.

The temperature history 22 is bounded by straight line segments. These temperature bounds indicate an increasing tempera maximum process temperature range for the remainder of the process.

The theoretical maximum density "history" 24 follow paths to maximum values which are represented by a generalized graph 24. The theoretical maximum density is defined as the closest possible packing of atoms into the crystalline structure of the compound; exclusive of any and all impurities, that will produce a minimum interstitial volume between the packed atoms.

The break away points as a function of time 30 vary, moreover, with the material and billet size under consideration.

An example, not in accordance with the invention, but useful for describing some features of the invention, will now be given.

Alpha alumina powder of less than one micron, preferably less than one tenth of a micron, particle size is worked or ball milled in a dry mill from four to eight hours. Preferably, alumina sold by W. R. Grace Company of U.S.A. under the name "Grace-KA 210" should be used as the raw material. This alumina powder has a surface area of approximately 9 meters2/gram. s} moreover, of very high purity, although it does contain 0.1% addition of MgO. Other aluminas also can be used, although experimental data does seem to indicate that best results are achieved with the Grace-KA 210 material.

To maintain powder purity, moreover, the ball mill also, should be formed from very pure alumina.

Upon completion of the milling step, the powder is baked for another four to eight hours at 50° to 100°C. Baking the powder at 72°C seems to' be a preferred temperature for this step in the process. These ball milling and drying operations appear to have the effect of removing excess surface gases to produce a finer-grained end product. The relation between the surface gas and the established. It is possible, however, that the surface gas behaves as an impurity phase that causes severe selective grain growth at high temperatures.

After outgassing, to produce a one-inch diameter billet of ^Oj, the powder is screened through a 200 mesh United States Standard sieve to break up any agglomerates that may have formed. The sifted powder is placed in a high temperature, high strength die. Typically, a graphite die in an inert, vacuum or reducing atmosphere is suitable for the purpose. A compacting pressure of 4000 to 8000 pounds per square inch, (psi) is applied to the powder within the die. This pressure is applied to initially compact the powder to 30% to 50% of its maximum theoretical density. For this sample, it has been found that an initial compacting or "prepressing" pressure of 5750 psi leads to the best end product results. This prepress force is then reduced to a range of 500 to 1000 psi.

Generally, a reduction in pressure to 1000 psi will produce acceptable results.

The powder and the die are placed in a hot press or other high temperature and high pressure sintering device. A protective atmosphere, moreover, is established in this system in order to preserve the die. A vacuum, a helium or other inert atmosphere, or a mixed atmosphere of inert gas and 8% by weight of hydrogen have been found suitable for this purpose. Furthermore, relatively less expensive nitrogen gas may be used for process economy.

Starting then with the reduced pressure on the compacted powder, the temperature of the powder and die is raised by means of an induction heater at a rate that is bounded by 400 to 1000°C per minute. By proper positioning and sizing of the induction heater and the billet generally uniform heating throughout the powder can be established. Within the above range it appears that manner until the onset of shrinkage of "break away point" 13 (Figure 1) is reached without degrading the quality of the final product.

With respect to the sample under consideration, numerous tests indicate that raising the temperature, within the above rate boundaries, of the powder and the die to 760° to 815° as measured with an optical pyrometer will produce the desired result. That is, the onset of shrinkage or "break away point" usually commences as the temperature reaches about 800°C. Preferably, while the temperature is being raised to the illustrative 800°C, to commence shrinkage, the reduced pressure of 1000 psi also is applied to the powder billet. This shrinkage may be observed with the aid of a linear variable displacement transducer that is attached to the ram that applies the pressure to the sintering powder.

After the "break away point" is reached, both temperature and pressure are increased in order to promote the rate of densi- fication that is inherent or natural to the particular material and billet size. Both pressure and temperature increase rates can be monitored and adjusted to approximate this natural rate.

With respect to the above alumina example, the As the application of this pressure continues, the temperature also is increased, but at a lower rate than that which characterized the initial increase to 800°C." Best results seem to be achieved with a maximum process temperature of about 1600°C that an optical pyrometer. This maximum temperature and pressure are s V sustained for two to six minutes, and preferably for three minutes, if a maximum process temperature of 1600°C is achieved. During this time, the alumina is sintering at its "natural*1 or inherent rate as referred to above.

The linear change in ram displacement between the times t£ and t3 shown in Figure 1 is a characteristic feature of a billet that is sintering at this natural rate. Other natural sintering rate indices are possible, although ram displacement is a most convenient technique.

The pressure and temperature increase rates that are applied to the sintering billet after the "break away point" 13 has been reached are generally adjusted to establish and maintain this natural sintering rate. The natural sintering rate will, of course, vary according to the material that is being processed. This natural rate, also may vary for different batches of the same . material. Consequently, the precise rates of increase of temperature and pressure that should be applied to the sintering billet for any particular material can be determined through a number of tests each performed on a different batch of the material. These tests will identify those conditions that produce the linear ram displacement 14 (Figure 1) , or other indications of the natural sintering rate, for the material under consideration. Once these sintering conditions are identified, subsequent billets can be processed without ram displacement observations and the like.

A more detailed consideration of Figure 1 indicates that the ram displacement is not entirely linear towards the completion of sintering 17. Thus, as shown in the drawing, the rate of ram displacement as a function of time decreases as the sintering billet approaches a condition of maximum densification. As this pressure and temperature applied to the billet is stabilized for two to six minutes to "cure" the now sintered billet.

Care must be excercised to terminate production conditions at this point in order to prevent the development of a "bloated" billet.' This "bloating" 16 is characterized by a reduced density billet, as indicated through the greater billet volume which the increasing ram displacement registers.

Turning once more to the completion of sintering 17, it is possible more precisely to promote the natural rate of sintering, which apparently changes as maximum densification is approached, by adjusting the temperature and pressure increase rate that is applied to the sintering billet in a manner that will enable the ram displacement to more nearly approximate the curve illustrated in Figure 1.

After the period of curing, or sustained heating at the maximum process temperature and pressure, the induction heater, or other source of heat, is turned off and the pressure on the alumina within the die is reduced to zero. A cooling period of one to five minutes is sufficient to enable the die (and the now sintered alumina) to cool to room temperature for removal from the press and separation from the die.

Samples of sintered alumina, produced in the foregoing manner, from the listed alumina powders, have shown in carefully executed laboratory tests the following characteristics: Avg. Khoop Number of Samples Hardness 1100* Grace-KA 100 8 2045 Grace-KA 210 21 2334 Other Commercial Sample A 10 2277 Avg. Compressive Avg. Modulus of Strength psi Rupture psi Grace-KA 100 326,700 44,100 Grace-KA 210 543,200 82,600 Other Commercial 321,000 ■ 59,500 Sample A Other Commercial 404,300 65,700 Sample B Modulus of Rupture Standard Deviation psi Grace-KA 100 16,500 Grace-KA 210 23,200 Other Commercial 16,000 Sample A Other Commercial 11,300 ' Sample B Compressive Strength Standard Deviation Grace-KA 100 115,000 Grace-KA 210 122,300 Other Commercial 111,600 Sample A Other Commercial 104,200 Sample B Average Grain Size Grace-KA 100 2.6 Grace-KA 210 0.72 Other Commercial Sample A 1.3 Other Commercial Sample B 1.7 material by means of a long, narrow, diamond shaped impression.

The hardness number is calculated as the ratio of the indenting r' load to the projected area of the indentation (THE MAKING, SHAPING AND TREATING OF STEEL, UNITED STATES STEEL 8th EDITION, 1964) .

In this connection it should be noted that the term "standard deviation" as used herein is the square root of the arithmetic mean of the squares of the deviations of the physical test data from their arithmetic mean.

The superior properties, on the average, of the sintered alumina that can be obtained if the Grace-KA 210 powder is used as a basic raw material in the described process is apparent. It should be noted that the Grace-KA 100 powder does not have an added 0.1% MgO crystal growth inhibitor. In developing the foregoing test data, moreover, sample preparation has been found to exert a significant influence. Chemical polishing of the samples, for instance, provides more realistic modulus of rupture test data. Mechanical polishing, however, seems to be detrimental to the actual strength of the sample that is undergoing testing.

Studies with a scanning electron microscope (at a magnification of 10,000) of the fracture surfaces of representative samples of alumina ceramic billets in the 1" to 5" diameter range that were produced in the manner described above demonstrate that the material has a grain size distribution as follows: Grain Size Range Percent of Grain Structure Less Than 0.3 micron 01 Between 0.3 and 0.5 micron 25% Between 0.3 and 0.7 micron 54% Between 0.3 and 0.9 micron 80% Between 0.3 and 1.5 micron 100% the relation between the diameter of the end product billet and the process conditions. Thus, to manufacture a five inch diameter billet of AI22, somewhat higher temperatures and pressures should be applied during processing than those conditions which are mentioned above with respect to the one inch diameter billet. It should be kept in mind, however, that a basic feature for all of the materials and billet sizes described herein is the application of an increased process pressure, within described boundaries throughout the sintering process, i.e. after the "break away point" (Figure 1). Moreover a maximum process pressure, an observed optimum, is identified within the described bounds, obtained by comparing- the pressure "history" of the sintering billet with the density of the processed sample, and may be more conveniently applied to the billet to provide the desired closest approach to the theoretical maximum density.

Thus, aluniina ceramics manufactured in accordance with the above-described method have a grain structure that is different from those grain sizes that have characterized earlier methods. Crystal of much larger average size, e.g. two or three microns generally occur in alumina made by these earlier methods.

This above-described method provides for the use of a metal oxide powder such as alumina powder, and a refractory carbide, nitride or oxide powder. Five inch diameter billets of alumina-titanium carbide (A^C^-TiC) were made from 70% by weight alumina powder (Grace-KA 210) and 30% by weight titanium carbide powder. The original particle size of the titanium carbide powder is 2 to 4 microns. The particle size is reduced by ball milling for 16 hours in alcohol, to an average particle size of 1 micron. The ball-milled powder is mechanically mixed with the alumina powder for uniform distribution of the two materials in the resulting TiC are blended together in an alcohol mixture in a ball mill for four hours. These mixed materials are removed from the ball mill, the alcohol is evaporated and the resulting powder is pre-pressed or compacted with physical pressure in the range of 400Q psi to 8000 psi to eliminate entrapped gas and to achieve a prepressed billet that has a density that is 30% to 50% of the maximum theoretical density. For the example under consideration, a 6300 psi prepressing pressure affords a suitable balance between powder packing and the elimination of entrapped gases. The applied ram pressure is then reduced to a range of 500 to 1000 psi. While this lower pressure is being applied, the material is heated at a rate that is not less than 400°C per minute nor more than 1000°C per minute until the onset of shrinkage commences, usually at about 800°C. While the material is being heated to this 800°C temperature the aforementioned reduced pressure is maintained to provide billet integrity, as noted above. With the onset of shrinkage that occurs at point 31 on the "break away" graph 30 in Figure 2, the ram pressure on the now sintering billet is increased to '5000 psi, - the preferred maximum hot process pressure. Suitable results, however, can be obtained with applied ram pressures in the 3000 to 9500 psi range.

As the application of this pressure continues, the temperature is increased, but at a lower rate than that which characterized the initial increase to 800°C. Thus, within six to ten ndnutes , the niaximum process temperature is reached in the range from 800°C to 1800°C, preferably to 1800°C. Based on available experimental data, best results are achieved with a temperature of about 1500°C. This maximum temperature and 5000 psi pressure are sustained for two to six minutes, the pressure being released and the material bein used at the end of the eriod and prior to bloating.

As shown in Table I below, the resulting material is superior to chemically similar materials that are produced through prior art processes.

Twenty, five-inch diameter billets of alumina-titanium carbide were fabricated in accordance with this method to demonstrate process reproducibility and the superior physical characteristics of the product.

The resulting density data for all 20 billets, is shown in Table I. The average billet density was 4.257 g/cc -0.07%, whereas the prior art density for this material is 4.21 g/cc. The term average, as used herein, is -the quotient of the arithmetic sum of the data divided by the number of data values used in calculating the sum.

TABLE I BILLET DENSITIES BILLET NUMBER DENSITIES ' (g/cc) 63 4.249 64 4.259 65 4.254 66 4.257 67 4.256 68 4.260 · 69 4.260 70 4.256 71 4.254 72 4.258 73 4.258 74 4.260 75 4.262 4.257 TABLE I BILLET DENSITIES BILLET NUMBER DENSITY (gm/cc) 77 4.258 78 4.258 79 4.256 80 4.260 81 4.254 82 4.260 Average 4.257 gm/cc Standard Deviation 0.003 gm/cc • 1 . CO.07%) The billets were ground top and bottom on a Blanchard Registered Trademark) model No. 11 grinder and diced into 21 blanks, each 3/4" square and 5/16" thick. From each of the 20 billets, two of the 21 blanks were randomly selected for transverse rupture strength tests, (TRS). The two selected rupture test blanks were each sliced into three 1/4" by 3/4" by 5/16" parallelpipeds to provide a total of six rupture specimens for each billet. The specimens were surface ground on all sides for edge sharpness and size uniformity.

The individual specimens were tested for transverse rupture strength by a three-point loading. The transverse rupture strength (TRS) results of these tests are tabulated in Table II.

In Table II, the average TRS of the six rupture specimens taken from each billet is tabulated below along with the standard deviation for this data. The overall average (the average of the average of each group of six samples) and the standard deviation of this overall average was found to be 124,333 ί 11,542 psi.

TABLE II BILLET TRANSVERSE RUPTURE STRENGTHS BILLET # SAMPLE TRS (PSI) AVE. TRS (PSI) 63 1 114,379 TABLE II BILLET TRANSVERSE RUPTURE STRENGTHS BILLET # SAMPLE # TRS(PSI) AVE. TRS (PSI) 3 89,388 124,557 4 133,232 -18,750 5 139,993 6 124,453 64 1 139,002 2 150,663 3 146,399 146,187 4 137,252 ± 5,967 5 149,353 6 154,456 65 1 136,167 2 150,363 3 120,494 113,799 4 104,129 ±28,545 5 59,861 6 111,783 66 1 89,962 2 127,815 3 78,107. 117,929 4 132,966 ±24,591 5 142,334 6 136,392 67 1 119,570 2 124,848 3 87,309 97,015 4 70,860 ±20,937 5 104,966 TABLE II BILLET TRANSVERSE RUPTURE STRENGTHS TRSfPSI) AVE. TRS (PSI) 135,543 139,812 125,472 132,003 118,659 * 7,904 131,558 140,972 136,554 106,444 89,163 117,195 148,363 . ¼1,286 125,498 97,147 161,221 147,558 84,116 130,232 108,674 ±26,180 145,953 133,870 141,234 128,750 135,508 137,381 131,122. - 6»832 149,525 138 146 BILLET # SAi.CPLE # TItS ( PS I ) AVE. TUS (PSI) 72 1 119,675 2 126,684 3 68,657 108,767 4 142,746 - 27,271 5 86,072 6 No Test 73 1 144,735 2 146,156 3 132,574 137,011 4 151,390 - 12, 47 5 133,613 6 113,596 74 1 116,096 2 136,419 3 99,935 120,030 4 118,341 il4,520 5 140,984 6 108,405 75 1 149,731 2 132,475 3 143,868 143,487 4 151,635 - 6,244 5 141,792 6 141,422 6 1 121,548 BILLET * SAMPLE # TKS (PSl) AW.. TItS 3 79,607 121,824 4 139,120 ^19,871 5 123,104 6 133,589 132,863 141,150 107,048 130,709 111,856 ±15,711 147,159 144,178 137,132 118,460 113,052 120,840 108,541 ¾8,356 96,384 151,509 103,270 118,557 110,070 120,667 124,957 ill, 877 139,715 127,435 69,338 90,804 BILLET * SAMPLE * TRS (PSI) AVE. TRS (PSI) 4 110,484 ±31,824 5 151,916 6 142,611 81 1 140,372 2 141,461 3 136,273 130,407 4 138,052 -16, 908 5 93,083 6 133,200 82 1 141,398 2 78,340 3 139,642 117,042 4 88,848 ±25,437 5 115,036 6 138,936 The broken transverse rupture specimens were then mounted and polished for hardness testing. Rockwell^ hardness tests were discontinued when three of the Rockwell indentors were ruined after application to only five billets. In passing, it should be noted that a Rockwell test is a measure of hardness as manifested by the resistance of the material to the penetration of an indentor in response to the application of a known load. The subscript, A in this test, indicates the load and indentor type used in the test for this material (THE MAKING, SHAPING AND TREATING OF STEEL, UNITED STATES STEEL, 8th EDITION, 1964). Hardness tests, however, were performed on all twenty billets. The hardness data are tabulated in Table III.

Although the Rockwell^ tests are not conclusive due to the above mentioned breakage problem, the average of five data points indicates a 0.8 increase in the Rockwell^ hardness over the prior art. This 0.8 increase is a significant improvement over the prior art because increases of 0.1 are of practical importance in the industry, e.g. tools are graded by increases of 0.1 in Rockwell^ hardness.

TABLE III BILLET HARDNESS BILLET # R. HARDNESS KNOOP HARDNESS 63 93.75 3557 64 93.78 3557 65 93.83 3557 66 93.78 3557 67 93.95 3557 68 N.D. 3557 69 N.D. 3557 70 N.D. 3557 71 N.D. 3227 72 N.D. 3557 73 N.D. 3557 74 N.D. 3557 75 N.D. 3557 76 N.D. 3557 77 N.D. 3227 78 N.D. 3557 79 N.D. 2940 80 N.D. 3227 81 N.D. 3557 82 N.D. 3557 Average 93.82 Average 3477 Two of the six broken transverse rupture specimens from each billet were photographed at a ten power (10X) magnification (Sample A) for macrohomogeneity, i.e. visible differences in the color of the sample material under inspection. Only one specimen from all of the samples studied showed an inho ogeneity (a 0.4 mm equivalent diameter titanium carbide particle) as enumerated in Table IV below. The equivalent size of the inhomogeneities listed in Table IV, moreover, are defined as the average of the major and minor axis of the inhomogeneity.

TABLE IV BILLET MACRO-HOMOGENEITY NUMBER OF VISIBLE . EQUIVALENT DIFFERENCES IN SIZE OF BILLET # ' SAMPLE A# COLOUR' ' INHOMOGENEITY 63 1 0 2 · 0 64 1 0 2 0 65 1 0 2 0 66 1 0 2 0 67 1 0 2 0 68 1 0 2 0 69 1 0 2 0 70 1 0 2 0 71 1 0 2 0 72 1 0 2 0 73 1 0 2 0 74 1 0 NUMBER OF VISIBLE EQUIVALENT DIFFERENCES IN SIZE OF BILLET SAMPLE M COLOUR INHOMOGENEITY 75 1 1 0.4 mm. 2 0 76 1 0 2 0 77 1 0 2 0 78 1 0 2 0 79 1 0 2 0 80 1 0 2 0 81 1 0 2 0 82 1 0 2 0 Two of the broken transverse rupture specimens from the remaining samples of each billet, Sample B, were randomly selected for micro-homogeneity. These micro-homogeneity samples were polished and photomicrographed at 900X magnification. The results tabulated in Table V below indicate that the average largest titanium carbide agglomerate was 12 microns, and the average titanium carbide grain was 4.82 microns. It should be noted that agglomerates are combinations of two or more grains into one mass.

• TABLE V BILLET MICRO-HOMOGENEITY LARGEST TiC TiC AGGLOMERATES AGGLOMERATE OVER 10 SAMPLE EQUIVALENT EQUIVALENT LARGEST TiC ' ■ B# DIAMETER DIAMETER Grain 63 1 0 2 2 6 5.5/ 64 1 0 4 /" 2 J? 3 5 65 1 14/" 1 " LARGEST TiC TiC AGGLOMERATES AGGLOMERATE OVER 10 BILLET SAMPLE EQUIVALENT EQUIVALENT. LARGEST TiC B ih DIAMETE DL'-METER GRAIN 66 1 9 *■ 0 5/* 2 7 ^ 0 5 68 1 14 /" 1 4 2 20.5/" 2 4 " 69 1 12 1 5.5 - 2 12/" 1 5.5 /" 70 71 1 14/** 1 4 A 2 2 4 JM 72 1 9 0 4/" 2 14 1 4 75 1 11 2 4 /" 2 19/ 3 5.5/" 76 1 7 " 0 6 2 15 1 5/" 77 78 1 9 0 , 5/" 2 10. 0 6> " 80 1 13 2 3/*' 2 17 1 5/" AVERAGE 12.0 ^ 1.08 4.82 t/ A scanning electron microscope indicates that the alumina grain size of this material is approximately the same as the grain size (0.3-1.5/0 of the sintered alumina alone. The TiC grain size, however, is approximately which was the size of the ball milled titanium carbide powder.

As described, the process produces a significantly improved product in comparison with the prior art. The increased density of alumina-titanium carbide indicates that the applied rate controlled sintering technique immediately following the "break away" point maximizes the densification of the material relative to that which was heretofore obtainable. This process, moreover, is applicable to other powdered materials once the "break away" point is determined and the inherent or natural densification rate for the material in question is established. 44318/2 /

Claims (11)

1. A method for producing a refractory material f comprising the steps of ball milling a refractory carbide, nitride or oxide powder, mixing a sinterable metal oxide powder and said carbide, nitride or oxide powder, compressing said mixed powders to eliminate entrapped gas, heating said compressed powders at a first rate of 400° to lOOO'C per minute, applying a. constant physical pressure to said powders while the powders are being heated at said rate, applying an increased physical pressure to said powders upon reaching an onset of powder shrinkage to attain a maximum hot process pressure, heating said powders at a second rate lower than said first rate to attain a maximum temperature while applying said increased pressure, maintaining said maximum hot process pressure and said maximum temperature for two to six minutes, and releasing the pressuring and cooling the material at the end of said period and prior to bloating.

2. A method according to Claim 1, wherein the sinterable metal oxide powder is alumina powder.

3. A method according to Claim 1, wherein the ; material ball-milled is carbide powder.

4. A method according to Claims 1 to 3 for producing an alumina carbide material comprising the steps of ball milling a carbide powder, mixing alumina powder and said carbide powder, compressing said mixed powders, heating said compressed powders at a first rate of 400° to 1000°C per minute, applying a constant physical pressure to said powders while the powders are being heated at said rate, applying an increased physical pressure to said powders upon reaching an onset of powder shrinkage to attain a maximum hot process pressure, heating said powders at a lower rate than said first rate to attain a maximum 44318/2

5. A method according to any preceding clafo wherein said maximum temperature^is in the range of 800° to 1800°C.

6. A method according to any preceding claim wherein said maximum process pressure is in the range of 3000 to 9500 pounds per square Inch.

7. A method for producing an alumina-carbide material comprising the steps of ball milling a carbide powder in alcohol to an average particle size of 1 y, mechanically mixing alumina powder and said ball milledcarbide powder, compressing said powders by a physical pressure of 6300 psi, reducing said pressure to 1000 psi, heating said powders at a rate of 400 *.o 1000°C per minute, holding said reduced pressure constant during said heating until said powders reach a temperature of 800°C, increasing said pressure to a maximum in the range of 3000 to 9500 psi, increasing said heating at a lower second rate than said first rate to a temperature of 1500°C in about seven minutes, maintaining said maximum pressure during said heating, and holding said maximum pressure and said 1500°C temperature for two to six minutes.

8. An alumina-titanium carbide material comprising physical attribute characterized by an average transverse rupture strength of 124,333 ± 11,542 psi.

9. An alumina-titanium carbide material according to Claim 8, further comprising physical attributes characterized by an average Rockwell^ hardness of 93.82.

10. An alumina-titanium carbide material according to Claim 8 further comprising physical attributes characterized by an average Knoop Hardness of 3477.

11. An alumina-titanium carbide material according to Claim 8 further comprising physical attributes characterized by an average titanium carbide grain size of the order of i 1 μ and an alumina grain size of the order of 0.3 to 1.5 y. HE:mr