SILICON NITRIDE HAVING HIGH TEMPERATURE AND METHOD

Abstract: A Si<u3>uN<u4>u polycrystalline ceramic body consisting essentially of a first phase of Si<u3>uN<u4>u and a second intergranular phase comprising SiO<u2>u and a densifying additive, wherein the intergranular phase is substantially completely crystalline. The Si<u3>uN<u4>u body is prepared by sintering a body of particulate material which is free of glass forming impurities or sintering and heat treating a body of particulate material which contains glass forming impurities.

Description: SILICON NITRIDE HAVING HIGH TEMPERATURE AND METHOD TECHNICAL FIELD This invention relates to polycrystalline ceramic bodies of Si- j ., and more particularly relates to such bodies having a substantially crystalline intergranular phase, and also relates to methods for preparing such bodies. BACKGROUND ART Si 3 N. powder characterized by cation impurities of 0.1 weight percent or less, a morphology of predominantly, crystal¬ line alpha phase and/or amorphous phase and fine particle size (3 microns or less average particle size as determined by B. E. T.), When consolidated with an additive such as MgO or Y ? ° 3 an< ^ sintered, is known to enable production of poly- crystalline bodies approaching theoretical density. See U.S. Patent 4,073,845, issued to S. T. Buljan, et al. on February 14, 1978, and assigned to GTE Sylvania Incorporated. Such powders may be consolidated into dense bodies by either hot pressing at less severe temperature and pressure conditions than are necessary with some less pure and less reactive pow¬ ders (See U.S. Patent 3,830,652 issued to G. E. Gazza) or by cold pressing and sintering, which is not possible with some less pure and less reactive powders. In the fabrication of such polycrystalline bodies, up to 25 weight percent of yttrium oxide or a lanthanide rare earth oxide such as Ce0 2 is typically added as a sintering or densifying aid. While such modifying additives are thought to deleteriously affe maximum attainable high temperature strength and creep res tance, nevertheless such additives are thought essential t achieve highest densities, that is, densities approaching theorectical density, which are also critical to the achie ment of high temperature strength. Unfortunately, it has been found that certain dense Si^N. bodies containing densifying additives, while exhibi ing excellent mechanical strength at room temperature and elevated temperatures up to about 1,000 °C, undergo signif cant and sometimes drastic reductions in strength at tem¬ peratures above 1,000 °C and particularly above 1200 °C. ' DISCLOSURE OF INVENTION In accordance with the invention it has been discover that dense polycrystalline bodies of a first phase of Si 3 N grains and a second intergranular phase of Si^N., Si0 2 , intentionally added densifying aids and unintentional impu rities exhibit remarkable retention of mechanical strength at temperatures above 1000 °C up to 1600 °C when the inter granular phase is substantially completely crystalline. A used herein, the terms "crystalline" and "polycrystalline" describe a solid composed of atoms, ions or molecules arra in a pattern which is repetitive in three dimensions and exhibits distinct x-ray diffraction intensity peaks charac istic of the crystal structure, while "glass", "glassy" or "amorphous" means a shor range order solid as indicated by its lack of x-ray diffraction intensity peaks and by broad infrared absorption peaks. In accordance with preferred embodiments of the inven tion to be described in detail herein, the crystallinity o the intergranular phase by the exclusion of amorphous or glassy material may be achieved either b -controlling the composition during processing to insure the exclusion of glass stabilizing additives or impurities, or in the alter native by giving the body a post-sintering crystallization heat treatment. Such methods enable production of polycrystalline bodie having improved mechanical strength at temperatures above 1000 °C up to 1600 °C, making such bodies useful as structu¬ ral parts or components in high operating temperature envi- ronments, for example, vehicular and aerospace engines and structures, regenerated for industrial waste heat recovery, etc. BRIEF DESCRIPTION OF THE DRAWING Figure 1 is a portion of an x-ray diffraction pattern f a silicon nitride body containing six percent Y 2 ° 3 P lus abou 2.5 percent Al^O^; Figure 2 is a portion of an x-ray diffraction pattern for the silicon nitride body of Figure 1 essentially without A1 2 0 3 ; Figures 3 and 4 are load deformation curves for Si 3 . bodies containing 12 percent Y 2 °3 Pl us U P to 2 weight percen Al 2 0 3 before and after a post-sintering crystallization heat treatment; Figure 5 is a portion of an x-ray diffraction pattern for a 52 percent Y ° 3 ~ 28 percent 1 2 0 3 - 20 percent Si0 2 composition; Figure 6 is a portion of an x-ray diffraction pattern for a 65 percent Y 2 °3 ~ ^ percent Si0 2 composition; Figure 7 is a dilato etric plot of curves "a" and "b" for the compositions of Fig. 5 before and after heat treat¬ ment; and Figure 8 is a plot of mechanical stresses at 1400 °C versus cross head travel for Si 3 N 4 bodies containing 6 per¬ cent ϊ 2 °3 and ° to 2 * 5 Percent 1 2 0 3 . BEST MODE FOR CARRYING OUT THE INVENTION For a better understanding of the present invention, together with other and further objects, advantages and capabilities thereof, reference is made to the following disclosure and appended claims in connection with the above- described drawings. The materials described herein have a composite micro- structure of Si 3 ^ grains and an intergranular phase consis¬ ting of Si0 2 (a natural impurity on the surface of the Si 3 N 4 starting powder) and the densifying additive. In addition, unintentional impurites present in the body tend to concen- trate in the intergranular phase. Because the intergranular phase is continuous, its characteristics profoundly affect the high temperature properties of the composite material. An amorphous second phase is highly undesirable in that sili- cate glasses which tend to form between the grains lead to poor mechanical strength of the Si 3 , body above about 1200 ° The presence of amorphous material between the Si 3 N 4 grains facilitates their separation at high temperature. This separation is manifested by slow crack propagation leading to plastic deformation of the body and causing degradation of structural integrity. Impurities in the starting raw material powders, or impurities introduced during processing prior to consolidation which are present in Si 3 4 bodies such as oxides (such as Al 2 0 3 ), which are soluble in the silicate second phase will stabilize this intergranular phase in its amorphous state by offering hindrance to the reorganization process which allows crystallization. It is therefore desirable for these powders to be processed into polycrystalline bodies in such a way that the intergranular phase undergoes full crystalli- zation during cooling from the sintering temperatures or alternatively that this full crystallization be obtained by a subsequent heat treatment procedure. By either course, a considerably more refractory intergranular phase is obtained which in turn leads to significantly enhanced high tempera- ture mechanical strength and a large reduction or elimination of the plastic deformation observed in the presence of glassy phases. While the invention applies to bodies containing any densifying additives, (known additives including for example, , . . - MgO, CrN, Zr0 2 , ZrN, Y 2 0 3 , La 2 0 3 and Ce0 2 ), it is preferred to employ a densifying additive of the type RE 2 0 3 where RE is yttrium or a lanthanide rare earth including La, Ce, Pr, Nd, Pm, Sm, Eu, Gd, Tb, Dy, Ho, Er, T , Yb, and Lu. While some of these rare earth oxides may not be in the form of RE 2 0 3 as starting materials, for example, Ce0 2 , and Pr0 2 , they are subsequently converted to the RE 2 0 3 form in the final dense polycrystalline body. To prepare powder compositions for consolidation, Si 3 N. powders are typically mixed with a densification aid by ball milling. The milling may be done wet and/or dry. There is a natural contamination of the powder mixture by abrasion of the milling container and milling media surfaces during milling. In order to substantially reduce or avoid such con- tamination, milling should be carried out in non-abradable or non-contaminating containers such as linear polyethlene, polyurethane or rubber lined containers using Si 3 N. milling media. Chemical analysis of powders milled both wet and dry with A1 2 0 3 media shows a total of 1 to 3 weight percent A1 2 0 3 added to the batch with about 1/2 of that amount picked up during wet milling and the remainder half during dry milling. This amount has been reduced by incorporation of non-contami¬ nating Si 3 N 4 media into the milling procedure and/or by shortening milling time. Using Si 3 N. media for wet milling and l 2 0 3 media for dry milling reduces the 1 2 0 3 content to about 0.5 to 1.5 weight percent. Milling both wet and dry with Si 3 N 4 media reduces the 1 2 0 3 content essentially to zero, that is, to the level present as an impurity in the Si 3 N 4 powder. The milled powder is usually screened through a coarse mesh screen, e.g., 50 or 100 mesh, prior to consolidation. Bodies may be prepared by hot pressing, hot isostatic pressing, or by any pressureless sintering step preceded by a suitable consolidation step such as dry pressing, isostatic -BUREAU OMPI pressing extruding, slip casting, injection molding, etc. See U.S. patent 4,073,845 for a general procedure for pres sureless sintering of Si 3 N. bodies. The screened powder i then loaded into a graphite hot pressing die whose interio surfaces have previously been coated with boron nitride po der. The powder is then prepressed at about 2000 psi and then the die is placed in a chamber containing argon and a pressure of about 500 psi is applied up to about 1200 °C, and then pressure and temperature are increased simulta- neously so that the ultimate pressure and temperature are achieved at about the same time. The densification proces is monitored using a dial guage indicating ram travel with the die body. A rate of downward movement of the ram cros head below about 0.004 inches per hour indicates completio of densification. The assembly is then cooled over a peri of about 1 to 2 hours. Ultimate pressures and temperature of from about 3,000 to 5,000 psi and 1675 °C to 1800 °C fo a time of about 2 to 5 hours are adequate conditions for t achievement of essentially full densification of the silic nitride body. The strength of compositions hot pressed to theoretical density from Si 3 .-Y 2 0 3 and Si 3 4 ~Ce0 2 powders with a range of 1 2 0 3 contents are given in Tables I and I All strength measurements are in terms of modulus of 3 rupture (MOR) (10 psi), and were made from room temperatur to 1400 °C using a four point loading fixture with a 0.9" outer span and a 0.4" inner span, which stresses the speci men by cross bending. Specimen size was .050" by 0.100" b one inch. All hot pressed specimens were tested in the strong direction with a cross head speed of .02" per minut on an Instron machine. Strength was determined by the ela stic beam formula, the dimensions of the test specimen and the test fixture. Theoretical density was calculated by the simple rule of mixtures for Si 3 N 4 and RE 2 0 3 . As may b seen from these Tables, strength increases dramatically at 1400 °C with decreasing Al 2 0 3 impurity content. TABLE I Strength of Hot Pressed Si 3 4 +13 Weight Percent Y 2 °3 MOR (psi X 10 ) Sample A1 2 0 3 weight Room No. percent Temperature 100 °C 1200 °C 1400°C 1 0 122.9 112.7 88.8 93.8 2 0.5 128.8 93.7 47.5 3 2.44 133.8 110.1 85.6 30.7 TABLE II Strength of Hot Pressed Si 3 4 ~Ce0 2 MOR (psi X IP" 3 ) 1.5 A1 2 0 3 No A1 2 0 3 (weight percent) Sample Ce0 2 weight Room Tem¬ Room Tem¬ No. percent perature 1400 °C perature .1400 °C 1 4 104.5 27.7 96.6 43.4 2 8 107.5 23.9 103.8 39. 3 12 102.9 17.8 134.3 75. The data presented in Tables I and II was obtained for module fabricated from high purity Si.,N 4 powder, that is, powder containing less than about 0.1 weight percent cation impurities. Tables III and IV show the effect of impurity content in the starting powder material on final high tem¬ perature strength. As may be seen from the Tables, when both powders are milled with Si-,N 4 media and in linear poly¬ ethylene containers so that no alumina is introduced during milling, high temperature strength is up to 60 percent greater for the high purity powder than for the low purity powder. Both samples were hot pressed at 1750 °C at 5000 psi to 100 percent of theoretical density. OMPI TABLE III Impurity Content of Starting Si 3 N 4 Powder (percent by weig Impurity Cation High Purity Powder Low Purity Powder Al 0.002 0.424 Fe - 1.15 Mn - 0.027 _ C - 0.458 Mg 0.0007 0.013 Ca 0.0006 0.224 Mo 0.01 TABLE IV Strength of Hot Pressed Si 3 N 4 tl3 Weight Precent Y 2 0 3 Comp'osition MORE (psi X 10 3 ) Si A Room - Starting Powder Temperature 100 °C 1200 °C 1400 °C High purity Powder 122.9 112.7 88.8 93.8 Low Purity Powder 98.9 88.3 65.7 57.9 The previous examples are for sintering by hot pressi Si 3 N 4 bodies. Pressureless sintering in a nitrogen atmos¬ phere was also used to prepare Si-N 4 bodies with varying A1 2 0 3 content. An Si.,N 4 body with 6 weight percent Y 2 0 3 a essentially no A1 2 0 3 was sintered two hours at 1775 °C and two hours at 1975 °C resulting in about 98 percent of theo retical density. This sample was subjected to x-ray diffr tion analysis using a scanning x-ray diffractometer with copper K radiation. Diffraction analysis showed three ph present, bets Si 3 N 4 , Y 2 Si 2 0 7 and 10Y 2 0 3 . 9Si0 2 . lSi 3 N 4 . Previous diffraction analysis of an Si 3 N 4 body containing six weight percent Y 2 0 3 and about 2.5 weight percent A1 2 0 3 OMP ■ WIP showed only beta Si 3 N 4 to be present, indicating that the second phase was x-ray amorphous. The key portions of the diffractometer scans are shown in Figures 1 and 2 for the alumina-containing and alumina-free bodies respectively, indicating that A1 2 0 3 causes the second phase to be x-ray amorphous in pressureless sintered Si, 4 bodies. The effect of varying A1 2 0 3 levels upon high temperature strength of pressureless sintered bodies of Si 3 N 4 containing 6 percent Y 2 °3 * s shown * n Table V and Figure 8, which is a plot of stress at 1400 °C versus cross head travel in 10 inches, indicating deformation of the sample during loading. The Figure graphically illustrates for the data of Table V that strength at 1400 °C is dramatically reduced by the presence of A1 2 0 3 in amounts up to 2.5 weight percent, where the intergranular phase has not " been heat treated to promote crystallization. TABLE V Strength of Pressureless Sintered Si 3 N 4 Bodies Containing 6 Percent. Y 2 0 3 3 MOR (psi x 10 J ) Sample Density (%) A1 2 0 3 Room No. Theoretical (wt. %) Temp. 1000°C 120°C 1400°C 1 98.0 0 83.8 68.8 57.1 2 98.2 0 69.4 62.2 64.4 61.3 3 100.0 1.5 100.6 - 65.5 29.5 4 100.0 2.5 104.7 88.6 56.7 23.4 A series of Si 3 N 4 + 12 weight percent Y 2 0 3 specimens con¬ taining about 1/2 to 2 weight percent A1 2 0 3 were- prepared and some were heat treated for 5 hours at 1525 °C to crystallize the 10-9-1 second phase after sintering. All specimens were tested to determine their strength. The recorder charts for as-sintered and for heat-treated specimens are given in Figure 3 and 4 where the applied load is horizontal and the cross- head travel time is vertical. The top two curves for as-sintered samples (Figure 3) showed rounded displacement (cross-head travel) curves as the load is applied with 50 t 80 seconds required to apply sufficient load to fracture the specimens. This curvature is indicative of plastic de- formation of the samples during loading. The average break ing strength of the two example specimens was 19,205 psi. For the two heat-treated samples, with the crystallized 10-9-1 second phase a different result was obtained as evi¬ denced by the load-deformation curves in Figure 4. First, there is essentially no curvature to the displacement trace indicating that plastic deformation has been minimized or eliminated. Only 20 to 22 seconds was needed to apply the breaking load, a further indication of the elimination of plastic deformation. Finally, the average breaking strengt for the two specimens was " 25,255 psi, an increase of 31.5% above that measured for the as-sintered specimens. This shows that a crystallized second phase formed by heat-treat will increase the elevated temperature of Si 3 N 4 -Si0 2 ~Y 2 0 3 compositions relative to the same composition with amorphou second phases in the as-sintered condition. In the above pressureless sintered samples, the second phase volume was small. In order to more distinctly observ the phase morphology of the intergranular material, two pow der mixtures of 52 weight percent 2 0 3 - 28 weight percent 1 2 0 3 - 20 weight percent Si0 2 ; and 65 weight percent Y 0 3 35 weight percent Si0 2 , respectively were prepared. These are the second phase compositions of the Si 3 N 4 body contain ing 6 weight percent Y 2 ° 3 w ith 2 weight percent A1 2 0 3 and without A1 2 0 3 , respectively. The Si0 2 was added since it i a natural species on the surface of the starting Si.,N. powd at about 3 weight percent level. The mixtures were melted at 1750° and 1840° respectively and then quickly cooled to room temperature. As melted, the Y 2 ° ~ Al j O., ~ s iO ? com¬ position contained the phase ^ 2 °3 ' 9si0 2' a nonequilibriu phase according to R.R. Wills et al., J. Materials Science, O OMM Vol. 11, pp. 1305-1309, 1976, plus a large amount of amorphous material as evidenced by the broad hump in the x-ray diffrac¬ tion scan of Figure 5. The non-Al 2 0 3 containing Y 2 0 ~ ^^2 composition crystallized directly to the equilibrium phase Y 2 Si 2 0 7 , according to Wills, et al., plus a minor amount of YgSig0 2 - j , with no evidence of amorphous material. See the x-ray diffraction scan of Figure 6. Thus, by direct experi¬ mentation with simulated intergranular compositions, it was demonstrated that the presence of a glass stabilizing oxide such as A1 2 0 3 hinders the crystallization process during cooling which otherwise would proceed spontaneously to a high degree of completion. In order to determine if the high temperature mechanical properties of the glassy intergranular phase contribute to the high temperature plastic deformation observed in A1 2 0 3 ~ containing Si 3 N 4 bodies, dilatometer tests were carried out on the intergranular phase simulated compositions prepared above. The curves are shown in Figure 7. Curve "a" indicates that the Al 2 0 3 - containing composition has a glass transi- tion of about 800 °C and a dilatometric softening point of about 890 °C. The transition temperature is that at which the thermal expansion changes from a relatively low value of a solid glass to the relatively high value of a liquid phase, while the softening point is the temperature at which the pressure of the dilatometer probe causes deformation of the sample. For comparison, curve "b" for the composition after a crystallization treatment at 1400 °C shows some evidence of a glass transition temperature at as high as 1200 °C shows some evidence of a glass transition temperature at as high as 1200 °C and a softening point at about 1380 β C. Thin indicates that the glass phase crystallized at about 1400 °C. Strength was measured for the Y 2 0 3 -Al 2 o 3 -si0 2 body as melted and also after heat treatment at 1400 °C for 15 hours. The samples were cyclically loaded to 6000 psi at successively higher temperatures to determine the non-elastic yield point. The unheat treated sampled was elastic to 800 and deformed at 900 °C. In contrast, the heat treated samp was elastic up to 1300 °C, and failed with no evidence of yielding at 1360 °C, indicating an MOR of 19.2 Ksi. Thus, without heat treatment both the dilatometer and MOR tests indicated a substantial decrease in viscosity of the as-mel sample at about 900 °C. X-ray diffraction analysis of the heat treated sample showed that heat treatment converted th non-equilibrium 7Y 2 0 3 . 9Si0 2 phase to Y 2 Si 2 0 7 and an un- identified crystalline phase with no eivdence of amorphous material. Confirmation of Y 2 Si 2 0 7 as the major phase was accomplished metallographically. INDUSTRIAL APPLICABILITY The Si 3 N 4 bodies of the present invention are useful a engine parts and components or regenerator or recuperator structures for waste heat recovery. OM

Claims: CLAIMS 1. A polycrystalline ceramic body consisting essen¬ tially of a first phase of Si-N 4 grains, and a second inter¬ granular phase comprising Si0 2 and a densifying additive, characterized in that the intergranular phase is substantially completely crystalline. 2. The body of Claim 1 wherein the densifying additive is RE 2 0 3 where RE is selected from the group consisting of Y and the lanthanide rare earths La, Ce, Pr, Nd, Pm, Sm, Eu, Gd, Tb, Dy, Ho, Er, Tm, Yb, Lu. 3. The body of Claim 2 wherein the intergranular phase consists essentially of Si0 2 and RE 2 0 3 . 4. The body of Claim 3 wherein RE is selected from the group consisting of Y and Ce. 5. A method for producing the polycrystalline ceramic body according to Claims 1 through 4, the method comprising sintering a body of particulate material containing glass stabilizing impurities, characterized in that following sin¬ tering, the body is heat treated to promote substantially complete crystallization of the intergranular phase. 6. The Method of Claim 5 wherein sintering is by hot pressing. 7. The method of Claim 5 wherein sintering is by pressureless sintering. 8. A method for producing the polycrystalline body of Claims 1 through 4, the method comprising sintering a body of particulate material, characterized in that the particulate material is maintained substantially free of glass stabilizing impurities during preparation for sinter- ing* 9. The method of Claim 8 wherein glass stabilizing impurities maintained below about 0.5 weight percent. OMPI S . W1P0

Inventors: QUAIKENBUSH C, REID J, SMITH T
Applicants: GTE LABORATORIES INC, GTE LABOR INC
Application year: 1979
Main classification: C04B35
IPC codes: C04B35