MIATERIAL TENGE ENGMEERIM ELSEVIER Materials Science and Engineering A244(1998)11-21 Sol-gel synthesis of ceramic matrix composites E D. Rodeghiero, B C. Moore, B.S. Wolkenberg, M. Wuthenow, O. K. Tse, EP Giannelis Department of Materials Science and Engineering, Cornell Uninersity, Ithaca, NY 14853-1501, US.A Abstract Sol-gel techniques have been used to produce various high temperature ceramic matrix composites including Ni/a-Al,O3, e/a-Al,O3, Ni/ZrO2, SiC(whisker)/a-Al,O3, and SiC(platelet)/a-Al,O,, as well as chemically modified versions of some of these ystems. In all cases, the composites have displayed uniform microstructures with a high degree of dispersion between the matrix nd reinforcement phases, a goal often not achieved when utilizing conventional powder mixing and processing techniques. The metal-ceramic composites investigated exhibit enhanced toughness and machinability as well as the potential for catalytic applications due to their novel fine-scale microstructure. Likewise, the Sic-reinforced alumina materials have been shown to be lighter, stiffer and tougher than pure alumina, without the use of the extreme hot-pressing temperatures and pressures needed by conventional powder processing approaches to produce the same results. o 1998 Elsevier Science S.A. All rights reserved Keywords: Sol-gel; Ceramic matrix composites; Microstructure 1. ntroduction processes, and as a result, virtually all sol-gel research since the late 1980s has been carried out in the thin Sol-gel processing received extensive attention in the film/coating area. However, this ignores sol-gels po- 1980s as literally hundreds of re- tential to play a supporting role in the synthesis of searchers sought after novel, low temperature methods monolithic structural ceramics and ceramic composites of producing common oxide ceramics such as silica, In other words, while the production of structural alumina, zirconia and titania in fully dense monolithic ceramics will most likely never be accomplished solely form [1]. Much of this excitement resulted from the by low temperature sol-gel techniques, the incorpora roduction of the first large-scale xerogels by Yoldas in tion of some sol-gel aspects into a broader synthesis 975 [2-5]. These self-supporting monolithic alumina scheme can nevertheless be highly beneficial. In fact, gels were highly porous (60-70%), but nevertheless many pioneers of the sol-gel community have felt this suggested the potential for producing fully dense ce- way from the very beginning. For instance, Roy et al. ramic components at reduced temperatures in a near- described their original goal in the sol-gel field as net-shape fashion. Due to the inherent fracture achievement of homogeneity on the finest possible associated with the drying and consolidation of bulk scale he production of mono-phasic glasses and gels, however, it later became accepted that sol-gel mono-phasic ceramic powders and precursors [7]. In- would instead be limited to a much smaller realm of deed, it was Roy who first brought sol-gel science to applications, namely the production of thin films, wh broad attention in the ceramics industry in the 1950s because of their planar geometry were not susceptible and 1960s for exactly this reason [2] to the formidable cracking problem of monoliths More recently, it is the incorporation of sol-gel Fortunately for the sol-gel community, the great techniques into the synthesis of ceramic matrix com- explosion in the microelectronics industry came during posites which seems especially appealing. In 1981, Rice the same time period. This translated into a lar and Becher demonstrated that ZrO,Al,O3 ceramic-ce demand for both thin film materials and thin film ramic composites produced through sol-gel approaches were superior to their ball milled, powder-derived coun- Corresponding author. Tel: +1 607 2556684: fax: +1 607 terparts in overall fracture properties [8,9]. In fact, in 2552365 this work the first demonstration of a simultaneous 0921-5093/98/S19.00 0 1998 Elsevier Science S.A. All rights reserved PIS0921-5093(9700821-6
Materials Science and Engineering A244 (1998) 11–21 Sol–gel synthesis of ceramic matrix composites E.D. Rodeghiero, B.C. Moore, B.S. Wolkenberg, M. Wuthenow, O.K. Tse, E.P. Giannelis * Department of Materials Science and Engineering, Cornell Uni6ersity, Ithaca, NY 14853-1501, USA Abstract Sol–gel techniques have been used to produce various high temperature ceramic matrix composites including Ni/a-Al2O3, Fe/a-Al2O3, Ni/ZrO2, SiC(whisker)/a-Al2O3, and SiC(platelet)/a-Al2O3, as well as chemically modified versions of some of these systems. In all cases, the composites have displayed uniform microstructures with a high degree of dispersion between the matrix and reinforcement phases, a goal often not achieved when utilizing conventional powder mixing and processing techniques. The metal–ceramic composites investigated exhibit enhanced toughness and machinability as well as the potential for catalytic applications due to their novel fine-scale microstructure. Likewise, the SiC-reinforced alumina materials have been shown to be lighter, stiffer and tougher than pure alumina, without the use of the extreme hot-pressing temperatures and pressures needed by conventional powder processing approaches to produce the same results. © 1998 Elsevier Science S.A. All rights reserved. Keywords: Sol–gel; Ceramic matrix composites; Microstructure 1. Introduction Sol–gel processing received extensive attention in the 1970s and early 1980s as literally hundreds of researchers sought after novel, low temperature methods of producing common oxide ceramics such as silica, alumina, zirconia and titania in fully dense monolithic form [1]. Much of this excitement resulted from the production of the first large-scale xerogels by Yoldas in 1975 [2–5]. These self-supporting monolithic alumina gels were highly porous (60–70%), but nevertheless suggested the potential for producing fully dense ceramic components at reduced temperatures in a nearnet-shape fashion. Due to the inherent fracture associated with the drying and consolidation of bulk gels, however, it later became accepted that sol–gel would instead be limited to a much smaller realm of applications, namely the production of thin films, which because of their planar geometry were not susceptible to the formidable cracking problem of monoliths [6]. Fortunately for the sol–gel community, the great explosion in the microelectronics industry came during the same time period. This translated into a large demand for both thin film materials and thin film processes, and as a result, virtually all sol–gel research since the late 1980s has been carried out in the thin film/coating area. However, this ignores sol–gel’s potential to play a supporting role in the synthesis of monolithic structural ceramics and ceramic composites. In other words, while the production of structural ceramics will most likely never be accomplished solely by low temperature sol–gel techniques, the incorporation of some sol–gel aspects into a broader synthesis scheme can nevertheless be highly beneficial. In fact, many pioneers of the sol–gel community have felt this way from the very beginning. For instance, Roy et al. described their original goal in the sol–gel field as achievement of ‘homogeneity on the finest possible scale’ in the production of mono-phasic glasses and mono-phasic ceramic powders and precursors [7]. Indeed, it was Roy who first brought sol–gel science to broad attention in the ceramics industry in the 1950s and 1960s for exactly this reason [2]. More recently, it is the incorporation of sol–gel techniques into the synthesis of ceramic matrix composites which seems especially appealing. In 1981, Rice and Becher demonstrated that ZrO2/Al2O3 ceramic–ceramic composites produced through sol–gel approaches were superior to their ball milled, powder-derived counterparts in overall fracture properties [8,9]. In fact, in this work the first demonstration of a simultaneous * Corresponding author. Tel.: +1 607 2556684; fax: +1 607 2552365. 0921-5093/98/$19.00 © 1998 Elsevier Science S.A. All rights reserved. PII S0921-5093(97)008 21-6
12 D. Rodeghiero et al / Materials Science and Engineering 4244(1998)11-21 increase in both fracture strength and fracture tough ness for the ZrO -reinforced Al,O3 system was re- ported. Rice attributed this beneficial behavior to the extreme homogeneity of the sol-gel derived com osites. In 1984. Hoffman et al. also demonstrated the extensive composite homogeneity and dispersion that could be achieved through sol-gel approaches by syn and CdS/SiO, (where the AgCl and Cds phases were in s fine-scale photosensitive composites such as AgCl/SiO crystalline form), to phase-separated phosphate and mixed oxide glasses such as CrPO4/SiO2, CePO/Sio and Nd, SIO2(where both the matrix and minor phases were amorphous)[10, 1l]. At the same time, Roy et al. were also using these same synthesis procedures to produce the first sol-gel derived metal-ceramic com- osites(including Cu/Al,O3, Ni/Al,O3, Cu/ZrO2 and 2-theta( degrees) Cu/SiO2)[12]. In fact, Hoffman et al. and Roy et al. Fi of XRD patterns for the processing of a 20/80 vol% collectively performed the most extensive work on the Ni al-ceramic composite: (a) dried, unreduced powder, sol-gel synthesis of ceramic matrix composites to the blet drogen reduced powder, and(c)1400C hot-pressed present day, thoroughly investigating over 30 different hemical systems. However, in their efforts no attention was given to high temperature consolidation, mechani- stiffer and tougher than pure alumina, these composites cal properties or structural applications. Finally, sol were consolidated at lower temperatures and pressures gel techniques have even been successfully used to than would have been required had conventional pow- roduce multilayer ceramic-ceramic composites [13]. In der processing techniques been used this work, the phenomenon of Liesegang band forma tion was used to produce two-dimensional precipitated CuCrOa layers in silica gels. The thickness and spacing 2. Experimental of these layers were shown to be tailorable, and the bands were also shown to survive sintering tempera- 2. 1. Metal-ceramic composite synthesis tures as high as 1100C, indicating that high tempera ure anisotropic composites could be produced To produce the Ni/alumina and Ni/zirconia metal In this paper, we review our synthesizing a ceramic composites, first a 0.15 M absolute ethanol range of different high temperature ceramic matrix composites using sol-gel techniques. These composites vary from being metal-ceramic in nature(e.g. Ni/ a Al2O3, Fe/a-Al2O3, etc. ) to ceramic-ceramic (e.g. SiC articulate reinforced a-AL, O3). Throughout this work, it is the physical and mechanical properties of the g Sic composites, the composite microstructures and the property: microstructure: synthesis relationships which are the elements of primary interest. The advantages 3 gained from using the sol-gel type syntheses in place of conventional powder mixing and processing are numer ous. For instance. in the case of the metal-ceramic composites, extremely fine (often nanoscale)mi- crostructure with a high degree of dispersion between the metal and ceramic phases have been produced. As a result, these composites exhibit enhanced toughness and durability as well as a simultaneous potential for catal- ysis applications. Likewise, the efforts at producing 2-theta(degree ol-gel derived, Sic-reinforced alumina composites have resulted in materials with highly uniform and Fig. 2 ce of XRD patterns for th Alo posite:(a) dried, uncal homogeneous morphologies without the presence of cined (b)900C air calcined powder, and(c)1750.C hot- Sic agglomerates. Furthermore, while being lighter, pressed pellet
12 E.D. Rodeghiero et al. / Materials Science and Engineering A244 (1998) 11–21 increase in both fracture strength and fracture toughness for the ZrO2-reinforced Al2O3 system was reported. Rice attributed this beneficial behavior to the extreme homogeneity of the sol–gel derived composites. In 1984, Hoffman et al. also demonstrated the extensive composite homogeneity and dispersion that could be achieved through sol–gel approaches by synthesizing various types of di-phasic gels ranging from fine-scale photosensitive composites such as AgCl/SiO2 and CdS/SiO2 (where the AgCl and CdS phases were in crystalline form), to phase-separated phosphate and mixed oxide glasses such as CrPO4/SiO2, CePO4/SiO2 and Nd2O3/SiO2 (where both the matrix and minor phases were amorphous) [10,11]. At the same time, Roy et al. were also using these same synthesis procedures to produce the first sol–gel derived metal–ceramic composites (including Cu/Al2O3, Ni/Al2O3, Cu/ZrO2 and Cu/SiO2) [12]. In fact, Hoffman et al. and Roy et al. collectively performed the most extensive work on the sol–gel synthesis of ceramic matrix composites to the present day, thoroughly investigating over 30 different chemical systems. However, in their efforts no attention was given to high temperature consolidation, mechanical properties or structural applications. Finally, sol– gel techniques have even been successfully used to produce multilayer ceramic–ceramic composites [13]. In this work, the phenomenon of Liesegang band formation was used to produce two-dimensional precipitated CuCrO4 layers in silica gels. The thickness and spacing of these layers were shown to be tailorable, and the bands were also shown to survive sintering temperatures as high as 1100°C, indicating that high temperature anisotropic composites could be produced. In this paper, we review our work in synthesizing a range of different high temperature ceramic matrix composites using sol–gel techniques. These composites vary from being metal–ceramic in nature (e.g. Ni/aAl2O3, Fe/a-Al2O3, etc.) to ceramic–ceramic (e.g. SiC particulate reinforced a-Al2O3). Throughout this work, it is the physical and mechanical properties of the composites, the composite microstructures and the property:microstructure:synthesis relationships which are the elements of primary interest. The advantages gained from using the sol–gel type syntheses in place of conventional powder mixing and processing are numerous. For instance, in the case of the metal–ceramic composites, extremely fine (often nanoscale) microstructures with a high degree of dispersion between the metal and ceramic phases have been produced. As a result, these composites exhibit enhanced toughness and durability as well as a simultaneous potential for catalysis applications. Likewise, the efforts at producing sol–gel derived, SiC-reinforced alumina composites have resulted in materials with highly uniform and homogeneous morphologies without the presence of SiC agglomerates. Furthermore, while being lighter, Fig. 1. Sequence of XRD patterns for the processing of a 20/80 vol.% Ni/a-Al2O3 metal–ceramic composite; (a) dried, unreduced powder, (b) 1000°C hydrogen reduced powder, and (c) 1400°C hot-pressed pellet. stiffer and tougher than pure alumina, these composites were consolidated at lower temperatures and pressures than would have been required had conventional powder processing techniques been used. 2. Experimental 2.1. Metal–ceramic composite synthesis To produce the Ni/alumina and Ni/zirconia metal– ceramic composites, first a 0.15 M absolute ethanol Fig. 2. Sequence of XRD patterns for the processing of a 20/80 vol.% SiC(whisker)/a-Al2O3 ceramic–ceramic composite; (a) dried, uncalcined powder, (b) 900°C air calcined powder, and (c) 1750°C hotpressed pellet
E D. Rodeghiero et al/ Materials Science and Engineering 4244(1998)11-21 13 EHT=20. 00 kV Photo No. =28 Detector= Q Fig. 3. Backscattered electron SEM micrograph of the microstructure of a 5/95 vol. Ni/a-AL,O, composite; (Ni= light contrast, a-Al, O,= dark solution of either aluminum isopropoxide, Altwo the gels were vacuum dried and OCH(CH3)213(Aldrich Chemical), or zirconium isopro and to 231 groun 230 E esh powder form with the use of an oxide isopropanol complex, Zr[OCH(CH3)2]4 agate and pestle (CH,),CHOH(Aldrich Chemical), was prepared and Reduction of the gro heated until boiling. Next an aqueous 0. 2 M solution of formed by placing the powders in a quartz tube furnace nickel formate dihydrate, Ni(CHO,)2 2H,O(Johnson and heating under flowing 99.99% hydrogen(20 cm3 Matthey), was prepared and added to the ethanol solu- min-)for I h at a temperature of 1000-1100C. The tion. The addition of the aqueous metal salt solution to role of this heat treatment was 2-fold. First, the metal the alkoxide solution typically caused immediate gela- salt was decomposed to its metallic state(either Ni or tion of the ceramic precursor. Nevertheless, the gel Fe). Second, the condensation reaction in the ceramic formed was stirred vigorously for at least an additional phase was driven to completion, forcing the elimination 10 min at 70oC, in order to ensure complete inter-d of all excess water and hydroxyls persion of the metallic and ceramic precursors. Natu The metal-ceramic powders which resulted after re- rally, the amount of metal salt solution added depended duction were then uniaxially hot-pressed at 1400oC and pon the final metal-ceramic composite compositio 10 MPa for 3 h in 0.5 in inside diameter a-Al,O3 dies desired. The gel was then transferred to crystallization This hot-pressing was carried out under a reducing gas dishes for drying at 100C for 24 h. Finally, the dried mixture of Co and Co2 flowing at a total rate of 10 gels were ground with an agate mortar and pestle to a cm min-. The oxygen partial pressure typically em- powder size of 230 mesh. ployed was 10-12 atm. The role of this CO atmosphere In the case of the Fe/a-Al2O3 composites, a 500 ml, was not to promote further reduction but rather to 0. 1 M solution containing an appropriate ratio of alu- prevent the occurrence of any reoxidation. (A thorough minum nitrate nonahydrate, Al(NO3)3 9H,O(Aldrich review of the thermodynamics of the Ni/Alyo system in Chemical), and ferric nitrate nonahydrate, regard to both the 1000oC hydrogen reduction and the Fe(NO3)3. 9H,O(Aldrich Chemical), was first synthe- 1400C CO/CO2 hot-pressing has been previously re- sized. Then while stirring, aqueous 1 M NaOH was ported [14]. The thermodynamic aspects of the other slowly added until the ph of the nitrate mixture sur- composite systems investigated here are very similar. passed 7. The gel which resulted was next centrifuged at Following hot-pressing, the sintered composite pellets 7000 rpm for I h and twice washed with deionized were removed from the dies with the use of a circular water. Then, after repeating the centrifuge/washing step diamond blade saw. High quality, optically smooth
E.D. Rodeghiero et al. / Materials Science and Engineering A244 (1998) 11–21 13 Fig. 3. Backscattered electron SEM micrograph of the microstructure of a 5/95 vol.% Ni/a-Al2O3 composite; (Ni=light contrast, a-Al2O3=dark contrast). solution of either aluminum isopropoxide, Al [OCH(CH3)2]3 (Aldrich Chemical), or zirconium isopropoxide isopropanol complex, Zr[OCH(CH3)2]4 · (CH3)2CHOH (Aldrich Chemical), was prepared and heated until boiling. Next an aqueous 0.2 M solution of nickel formate dihydrate, Ni(CHO2)2 · 2H2O (Johnson Matthey), was prepared and added to the ethanol solution. The addition of the aqueous metal salt solution to the alkoxide solution typically caused immediate gelation of the ceramic precursor. Nevertheless, the gel formed was stirred vigorously for at least an additional 10 min at 70°C, in order to ensure complete inter-dispersion of the metallic and ceramic precursors. Naturally, the amount of metal salt solution added depended upon the final metal–ceramic composite composition desired. The gel was then transferred to crystallization dishes for drying at 100°C for 24 h. Finally, the dried gels were ground with an agate mortar and pestle to a powder size of 230 mesh. In the case of the Fe/a–Al2O3 composites, a 500 ml, 0.1 M solution containing an appropriate ratio of aluminum nitrate nonahydrate, Al(NO3)3 · 9H2O (Aldrich Chemical), and ferric nitrate nonahydrate, Fe(NO3)3 · 9H2O (Aldrich Chemical), was first synthesized. Then while stirring, aqueous 1 M NaOH was slowly added until the pH of the nitrate mixture surpassed 7. The gel which resulted was next centrifuged at 7000 rpm for 1 h and twice washed with deionized water. Then, after repeating the centrifuge/washing step two more times, the gels were vacuum dried and ground to 230 mesh powder form with the use of an agate mortar and pestle. Reduction of the ground precursor gels was performed by placing the powders in a quartz tube furnace and heating under flowing 99.99% hydrogen (20 cm3 min−1 ) for 1 h at a temperature of 1000–1100°C. The role of this heat treatment was 2-fold. First, the metal salt was decomposed to its metallic state (either Ni or Fe). Second, the condensation reaction in the ceramic phase was driven to completion, forcing the elimination of all excess water and hydroxyls. The metal–ceramic powders which resulted after reduction were then uniaxially hot-pressed at 1400°C and 10 MPa for 3 h in 0.5 in. inside diameter a-Al2O3 dies. This hot-pressing was carried out under a reducing gas mixture of CO and CO2 flowing at a total rate of 10 cm3 min−1 . The oxygen partial pressure typically employed was 10−12 atm. The role of this CO atmosphere was not to promote further reduction but rather to prevent the occurrence of any reoxidation. (A thorough review of the thermodynamics of the Ni/Al/O system in regard to both the 1000°C hydrogen reduction and the 1400°C CO/CO2 hot-pressing has been previously reported [14]. The thermodynamic aspects of the other composite systems investigated here are very similar.) Following hot-pressing, the sintered composite pellets were removed from the dies with the use of a circular diamond blade saw. High quality, optically smooth
E D. Rodeghiero et al/ Materials Science and Engineering 4244(1998)11-21 EHT=10.00 KV 8 mm 30um Photo No =18 Detector= Q Fig. 4. Backscattered electron SEM micrograph of the microstructure of a 50/50 vol. Fe/a-Al20 site: (Fe= light contrast, Al,,= dark surfaces were prepared by flattening the pellet faces 2.3. Ceramic-ceramic composite synthesis with a 20 um metal bonded diamond wheel(Struers), coarse polishing with Sic paper, and then fine polishing The preparation of the Sic(particulate)-reinforced with 6 um diamond paste impregnated Texmet polish lumina composites consisted of the following steps ing cloth(Buehler). Finally, an aqueous ultrasonic First a 0.15m absolute ethanol solution of aluminum to remove pellet surface contamination ier polishing isopropoxide was again prepared and heated to boil- cleaning bath was utilized immediately ing. Next, an appropriate amount of either SiC whiskers(≈ I um diameter by≈l5 um long) or Sic 2. 2. Doped metal-ceramic composite synthesi platelets(0.5-5 um thick by 5-70 um diameter)(John In certain instances. some of the metal-ceramic com- son Matthey) was added to the ceramic precursor solu- osites were doped with additional phases or com tion while stirring. After 5 min, just enough water to pounds. Two of the more extensively investigated gel the precursor solution was added. The mixture was xamples were Ni/a-Al2O3 doped with ZrO2 and Ni/a left to stir continuously at 70C until the gel was Al,O, doped with Cr2O3. In the case of the Zro viscous enough to prevent SiC settling. The gel was then transferred to crystallization dishes and dried for pared by simply adding small amounts of the zirconium 24 h at 100C. Finally, the dried gels were delicately alkoxide to the initial aluminum isopropoxide solution. hand ground and sieved in the same fashion as the Reduction and hot-pressing were then carried out as metal-ceramic precursors. To fully transform the ce- described above. To produce the Cr,O, doped material, ramic gel to A2 O, the dried powders were calcined in a chromium formate salt (added to the nickel formate a quartz tube furnace at 900 C in air for 1 h. This was dihydrate solution) was used as the dopant source. then followed by uniaxial hot-pressing at 1750.C and Again, reduction and hot-pressing were carried out 35 MPa for 3 h. The 0.5 in. inside diameter dies used normally. In this case, however, even though the in this case consisted of high strength graphite(Poco chromium source was incorporated along with the N raphite, grade ZxF-5Q) and were enclosed in a the form of a salt, the employed hydrogen reduction chamber backfilled with argon for the duration of the temperature of 1000 C was not high enough to reduce high temperature exposure. Handling and post-process- the chromium to its metallic state and hence, an alu- ing of the sintered Sic/a-Al2O3 was performed in a mina-rich Al,O3/Cr,O similar fashion as described for the metal-ceramic duced as the ceramic phase composites
14 E.D. Rodeghiero et al. / Materials Science and Engineering A244 (1998) 11–21 Fig. 4. Backscattered electron SEM micrograph of the microstructure of a 50/50 vol.% Fe/a-Al2O3 composite; (Fe=light contrast, a-Al2O3=dark contrast). surfaces were prepared by flattening the pellet faces with a 20 mm metal bonded diamond wheel (Struers), coarse polishing with SiC paper, and then fine polishing with 6 mm diamond paste impregnated Texmet® polishing cloth (Buehler). Finally, an aqueous ultrasonic cleaning bath was utilized immediately after polishing to remove pellet surface contamination. 2.2. Doped metal–ceramic composite synthesis In certain instances, some of the metal–ceramic composites were doped with additional phases or compounds. Two of the more extensively investigated examples were Ni/a-Al2O3 doped with ZrO2 and Ni/aAl2O3 doped with Cr2O3. In the case of the ZrO2 modified Ni/a-Al2O3, the doped composites were prepared by simply adding small amounts of the zirconium alkoxide to the initial aluminum isopropoxide solution. Reduction and hot-pressing were then carried out as described above. To produce the Cr2O3 doped material, a chromium formate salt (added to the nickel formate dihydrate solution) was used as the dopant source. Again, reduction and hot-pressing were carried out normally. In this case, however, even though the chromium source was incorporated along with the Ni in the form of a salt, the employed hydrogen reduction temperature of 1000°C was not high enough to reduce the chromium to its metallic state, and hence, an alumina-rich Al2O3/Cr2O3 solid solution (ruby) was produced as the ceramic phase. 2.3. Ceramic–ceramic composite synthesis The preparation of the SiC(particulate)-reinforced alumina composites consisted of the following steps. First, a 0.15 M absolute ethanol solution of aluminum isopropoxide was again prepared and heated to boiling. Next, an appropriate amount of either SiC whiskers (:1 mm diameter by :15 mm long) or SiC platelets (0.5–5 mm thick by 5–70 mm diameter) (Johnson Matthey) was added to the ceramic precursor solution while stirring. After 5 min, just enough water to gel the precursor solution was added. The mixture was left to stir continuously at 70°C until the gel was viscous enough to prevent SiC settling. The gel was then transferred to crystallization dishes and dried for 24 h at 100°C. Finally, the dried gels were delicately hand ground and sieved in the same fashion as the metal–ceramic precursors. To fully transform the ceramic gel to Al2O3, the dried powders were calcined in a quartz tube furnace at 900°C in air for 1 h. This was then followed by uniaxial hot-pressing at 1750°C and 35 MPa for 3 h. The 0.5 in. inside diameter dies used in this case consisted of high strength graphite (Poco Graphite, grade ZXF-5Q) and were enclosed in a chamber backfilled with argon for the duration of the high temperature exposure. Handling and post-processing of the sintered SiC/a-Al2O3 was performed in a similar fashion as described for the metal–ceramic composites
D. Rodeghiero et al/ Materials Science and Engineering 4244(1998)11-21 2. 4. Characterization The density of all hot-pressed composites was deter mined by calculating the mass and volume of each X-ray diffraction(XRD) of both the precursor and pellet after cutting into a prismatic configuration. In heat treated powders as well as the sintered pellets wa addition, the relative or percentage density of each performed using a Scintag Pad X diffractometer with material was calculated by dividing the measured den CuKa radiation. Reduced metal-ceramic powders were sity by the theoretical density based on composite com- encapsulated in 1.0 mm diameter glass capillaries under position. The resulting number, always slightly argon to prevent reoxidation of the very small metallic < 100%, was used as an indicator of residual porosity particles present at this stage of processing In the case of the metal-ceramic composites, electrical Polished, sintered microstructures were investigated resistivity of the hot-pressed pellets was also measured with both optical and scanning electron microscopy with the use of a standard two-probe ohmmeter, in (SEM). The optical microscope used was an inverted order to analyze percolation of the metallic phase. Only PME Olympus instrument while the scanning electron approximate measurements needed to be performed microscope employed was a Leica 440 Stereoscan ma- since the quantitative difference between percolated and chine with both secondary and backscattered electron non-percolated readings was several orders of magni- imaging capabilities. The SEM was also occasionally tude sed to image composite fracture surfaces The elastic constants of the isotropic metal-ceramic composites were determined using an ultrasonic tech- nique previously reported [15]. a similar approach was adopted for the SiC/a-Al2O3 materials; however, due to the uniaxial nature of the hot-pressing employed, the Sic whiskers and platelets had a tendency to lie perpen- dicular to the hot-pressing direction, resulting in an- isotropic composites of the transverse isotropy' type symmetry [16]. This required the incorporation of a quipment was needed but measurements had to be performed in several additional directions including at platelet alignment. The exact mathematics of this ex tended analysis will not be presented here, but the reader is referred to the derivations of Neighbours and Schacher for more insight [17] fracture toughness testing was performed by cutting 25m the sintered composites into beams, machining chevron notches into the beams, and then breaking the speci mens to complete failure on an Instron Model 1125 mechanical testing instrument equipped with a three point bending fixture. A linear variable displacement transducer(RDP-Electrosense, model RDP D5/10G8) was used to measure beam load-point displacement while a piezoelectric transducer (Kistler Instrument, model 9301A) recorded specimen load under constant displacement rate conditions. Data in the form of load/ displacement curves were collected with the use of a computerized data acquisition system. The beam di- mensions used throughout the testing were 1.9 x 1.9 8.0 mm. and the chevron notches were machined with a Well wire saw (Ahlburg Technical Equipment, model 3242) using 220 um diameter diamond impreg nated steel wire. The included angle of the chevron notch was maintained at x94o while all other sample and notch dimensions were in compliance with the Fig. 5. Optical micrographs of the microstructure of a 20 /80 vol% SiC(whisker)/a-Al,O, composite;(a) pellet face perpendicular to hot. work by Wu[18]. A much more thorough explanation ressing direction and(b) pellet face parallel to hot-pressing direction; of the toughness testing apparatus and process used is SiC=light contrast, a-Al, O,=dark contrast currently in preparation [191
E.D. Rodeghiero et al. / Materials Science and Engineering A244 (1998) 11–21 15 2.4. Characterization X-ray diffraction (XRD) of both the precursor and heat treated powders as well as the sintered pellets was performed using a Scintag Pad X diffractometer with CuKa radiation. Reduced metal–ceramic powders were encapsulated in 1.0 mm diameter glass capillaries under argon to prevent reoxidation of the very small metallic particles present at this stage of processing. Polished, sintered microstructures were investigated with both optical and scanning electron microscopy (SEM). The optical microscope used was an inverted PME Olympus instrument while the scanning electron microscope employed was a Leica 440 Stereoscan machine with both secondary and backscattered electron imaging capabilities. The SEM was also occasionally used to image composite fracture surfaces. The density of all hot-pressed composites was determined by calculating the mass and volume of each pellet after cutting into a prismatic configuration. In addition, the relative or percentage density of each material was calculated by dividing the measured density by the theoretical density based on composite composition. The resulting number, always slightly B100%, was used as an indicator of residual porosity. In the case of the metal–ceramic composites, electrical resistivity of the hot-pressed pellets was also measured, with the use of a standard two-probe ohmmeter, in order to analyze percolation of the metallic phase. Only approximate measurements needed to be performed since the quantitative difference between percolated and non-percolated readings was several orders of magnitude. The elastic constants of the isotropic metal–ceramic composites were determined using an ultrasonic technique previously reported [15]. A similar approach was adopted for the SiC/a-Al2O3 materials; however, due to the uniaxial nature of the hot-pressing employed, the SiC whiskers and platelets had a tendency to lie perpendicular to the hot-pressing direction, resulting in anisotropic composites of the ‘transverse isotropy’ type symmetry [16]. This required the incorporation of a much more complex acoustic analysis. (No additional equipment was needed but measurements had to be performed in several additional directions including at least one direction oblique to the plane of whisker/ platelet alignment.) The exact mathematics of this extended analysis will not be presented here, but the reader is referred to the derivations of Neighbours and Schacher for more insight [17]. Fracture toughness testing was performed by cutting the sintered composites into beams, machining chevron notches into the beams, and then breaking the specimens to complete failure on an Instron Model 1125 mechanical testing instrument equipped with a threepoint bending fixture. A linear variable displacement transducer (RDP-Electrosense, model RDP D5/10G8) was used to measure beam load-point displacement while a piezoelectric transducer (Kistler Instrument, model 9301A) recorded specimen load under constant displacement rate conditions. Data in the form of load/ displacement curves were collected with the use of a computerized data acquisition system. The beam dimensions used throughout the testing were 1.9×1.9× 8.0 mm3 , and the chevron notches were machined with a Well® wire saw (Ahlburg Technical Equipment, model 3242) using 220 mm diameter diamond impregnated steel wire. The included angle of the chevron notch was maintained at :94° while all other sample and notch dimensions were in compliance with the work by Wu [18]. A much more thorough explanation of the toughness testing apparatus and process used is currently in preparation [19]. Fig. 5. Optical micrographs of the microstructure of a 20/80 vol.% SiC(whisker)/a-Al2O3 composite; (a) pellet face perpendicular to hotpressing direction and (b) pellet face parallel to hot-pressing direction; (SiC=light contrast, a-Al2O3=dark contrast)