《复合材料 Composites》课程教学资源（学习资料）第五章 陶瓷基复合材料_j.1151-2916.2003.tb03272.x[1]

Au cern.Soc.8626-32(2003) ournal Preparation of SiC/SiC Composites by Hot Pressing, Using Tyranno-SA Fiber as Reinforcement Shaoming Dong, t Yutai Katoh, *and Akira Kohyama Institute of Advanced Energy, Kyoto University, Gokasho, Uji, Kyoto 611-0011, Japan The development of advanced Tyranno SA SiC fiber with a high temperature, because such treatment might cause creep defor near-stoichiometric composition and a well-crystallized micro- mation, leading to strength degradation, even though Hi-Nicalon is an structure has made it possible to prepare SiC/SiC composites free SiC fiber even under harsh conditions. To assess the reinforcing effec- With the development of advanced SiC-based fibers, well- iveness of Tyranno SA fiber at high temperature under crystallized and near-stoichiometric SiC-fiber compositions, such as pressure, unidirectional SiC/SiC composites were prepared Tyranno SA(UBE Industries, Ltd, Yamaguchi, Japan) fiber, have hot pressing, using pyrolytic carbon(PyC)-coated Tyranno been achieved. The newly available Si-Al-C Tyranno Sa fiber also fiber as a reinforcement and nanopowder Sic with sintering has a high tensile strength and elastic modulus and exhibits no additives for matrix formation. The effects of sintering condi- strength degradation or compositional changes under heating to tions on the microstructural evolution and mechanical prop- 1900C in an inert atmosphere or in air at 1000 C. The absence of erties of the composites were characterized. As the sinterin oxygen and perfect crystallization in this fiber make it a candidate for temperature increased(from 1720 to 1780C)and the sinter fabricating SiC/Sic composites at relatively high temperature and ing pressure increased (from 15 to 20 MPa), the density of the under harsh conditions, such as HP. However, challenges still exist for composites gradually increased. Simultaneously, the elastic e preparation of SiC/SiC composites by HP. The effectiveness of us, the proportional limit stress, and the strength, under fiber reinforcement under combined high temperature and pressure oth bend and tensile tests, also improved. At lower tempera- must be investigated ture and/or pressure, long fiber pullout was a predominant The purpose of this paper is to evaluate the effects of preparation fracture behavior, indicating relatively weak fiber/matrix conditions on the densification process and mechanical performance onding. However, at high temperature and/or pressure, short of unidirectional (UD) SiC/Sic composites fabricated by HP, via fiber pullout became a main fracture characteristic, indicating uid-phase sintering, using Tyranno SA as a reinforcement. Corre- relatively strong fiber/matrix bonding. These phenomena were lations among sintering temperature and pressure, microstructure, and also confirmed by the characteristics of the hysteresis loops physical and mechanical behaviors are discussed derived from the stress-strain curves produced by a tensile test present investigation, he reinforcement of Tyranno SA fiber is effective for provid IL. Experimental Procedure ng noncatastrophic fracture behavior to cor PyC-coated Tyranno SA fiber tows were used as a reinforce- iC/SiC composite fabrication in the present study ypical properties of the newly produced( Grade I) Tyranno SA . Introduction fiber are listed in Table I. The thickness of the fiber coating was 20.8 um. To lower the sintering temperature and promote N RECENT years, low-temperature fabrication techniques, such as densification, B-SiC nanopowder(Marketech International chemical vapor infiltration(CVi), polymer impregn Port Townsend, Wa) was used for matrix formation, and Al-O pyrolysis(PIP), and reaction sintering(RS) have been suc (Sumitomo Chemical Industries, Ltd, Tokyo, Japan)and Y,O applied for the fabrication of high-performance SiC/SiC (Johnson Matthey, Inc, Cheshire, U.K. were used as sintering ites. Composites with various chemical and mechanical aids. Typical compositions of the SiC nanopowder were have been developed by modifying interfacial properties(using SiC, 1%02% free carbon, and 1%1.5% oxygen, the ay fiber coatings) and infiltration techniques. particle size was <30 nm. Hot pressing(HP) is an effective processing technique for densi- Because the nanosized Sic powder used for matrix formation f ing powder compacts, especially when a second phase, such as has a very high specific surface area (110 m?/g), it was very fibers, is included. 2 Earlier developed SiC-based fibers, such ficult to Nicalon(Nippon Carbon Co., Tokyo, Japan)fiber, are thermodynam- winding method. In the present experiment, the PyC-coated ically unstable because of their high oxygen content; thus, high- tows were first wound and fixed onto frames, to form aligned UD temperature applications of these fibers for SiC/SiC composite fabri- fiber sheets. These sheets were then impregnated by a polymer cation are limited, because of decomposition of the Si-o-C phase. On containing slurry(polycarbosilane(PCS)mixed with a filler the other hand, even fibers with low crystallinity, such as Hi-Nicalon which was composed of SiC nanopowder and sintering additives (Nippon Carbon), cannot withstand high-pressure heat treatment at and pyrolyzed in an argon atmosphere for pretreatment. The polymer precursor of PCS(Nippon Carbon Co., Tokyo, Japan)was included to decrease the viscosity of the slurry and increase the wettability between the fibers and the matrix slurry. In this Icontributing editor xperiment, hexane was used as solvent. During polymer(with ller) impregnation and pyrolysis(PlP), the intrabundle matrix was compacted. The polymer-containing slurry needed to be carefully modified to increase the intrabundle infiltration effi- Manuscript No. 187345, Received November 7, 2001; approved October 10, 20 ciency and to avoid matrix cracking. Because of the very high pecific surface area(110 m/g) of nanopowder, higher SiC stitute of Ceramics, Chinese Academy of Sciences, addition would greatly increase the viscosity of the slurry at a hanghai 200050, People certain amount of solvent so that the infiltration efficiency was

Preparation of SiC/SiC Composites by Hot Pressing, Using Tyranno-SA Fiber as Reinforcement Shaoming Dong,† Yutai Katoh,* and Akira Kohyama* Institute of Advanced Energy, Kyoto University, Gokasho, Uji, Kyoto 611-0011, Japan The development of advanced Tyranno SA SiC fiber with a near-stoichiometric composition and a well-crystallized micro￾structure has made it possible to prepare SiC/SiC composites even under harsh conditions. To assess the reinforcing effec￾tiveness of Tyranno SA fiber at high temperature under pressure, unidirectional SiC/SiC composites were prepared by hot pressing, using pyrolytic carbon (PyC)-coated Tyranno SA fiber as a reinforcement and nanopowder SiC with sintering additives for matrix formation. The effects of sintering condi￾tions on the microstructural evolution and mechanical prop￾erties of the composites were characterized. As the sintering temperature increased (from 1720° to 1780°C) and the sinter￾ing pressure increased (from 15 to 20 MPa), the density of the composites gradually increased. Simultaneously, the elastic modulus, the proportional limit stress, and the strength, under both bend and tensile tests, also improved. At lower tempera￾ture and/or pressure, long fiber pullout was a predominant fracture behavior, indicating relatively weak fiber/matrix bonding. However, at high temperature and/or pressure, short fiber pullout became a main fracture characteristic, indicating relatively strong fiber/matrix bonding. These phenomena were also confirmed by the characteristics of the hysteresis loops derived from the stress–strain curves produced by a tensile test with unloading–reloading cycles. In the present investigation, the reinforcement of Tyranno SA fiber is effective for provid￾ing noncatastrophic fracture behavior to composites. I. Introduction I N RECENT years, low-temperature fabrication techniques, such as chemical vapor infiltration (CVI), polymer impregnation and pyrolysis (PIP), and reaction sintering (RS) have been successfully applied for the fabrication of high-performance SiC/SiC compos￾ites. Composites with various chemical and mechanical behaviors have been developed by modifying interfacial properties (using fiber coatings) and infiltration techniques. Hot pressing (HP) is an effective processing technique for densi￾fying powder compacts, especially when a second phase, such as fibers, is included.1,2 Earlier developed SiC-based fibers, such as Nicalon (Nippon Carbon Co., Tokyo, Japan) fiber, are thermodynam￾ically unstable because of their high oxygen content;3,4 thus, high￾temperature applications of these fibers for SiC/SiC composite fabri￾cation are limited, because of decomposition of the Si-O-C phase.5 On the other hand, even fibers with low crystallinity, such as Hi-Nicalon (Nippon Carbon), cannot withstand high-pressure heat treatment at high temperature, because such treatment might cause creep defor￾mation, leading to strength degradation, even though Hi-Nicalon is an oxygen-free SiC fiber.6 With the development of advanced SiC-based fibers, well￾crystallized and near-stoichiometric SiC-fiber compositions, such as Tyranno SA (UBE Industries, Ltd., Yamaguchi, Japan) fiber, have been achieved.7,8 The newly available Si-Al-C Tyranno SA fiber also has a high tensile strength and elastic modulus and exhibits no strength degradation or compositional changes under heating to 1900°C in an inert atmosphere or in air at 1000°C. The absence of oxygen and perfect crystallization in this fiber make it a candidate for fabricating SiC/SiC composites at relatively high temperature and under harsh conditions, such as HP. However, challenges still exist for the preparation of SiC/SiC composites by HP. The effectiveness of fiber reinforcement under combined high temperature and pressure must be investigated. The purpose of this paper is to evaluate the effects of preparation conditions on the densification process and mechanical performance of unidirectional (UD) SiC/SiC composites fabricated by HP, via liquid-phase sintering, using Tyranno SA as a reinforcement. Corre￾lations among sintering temperature and pressure, microstructure, and physical and mechanical behaviors are discussed. II. Experimental Procedure PyC-coated Tyranno SA fiber tows were used as a reinforce￾ment for UD SiC/SiC composite fabrication in the present study. Typical properties of the newly produced (Grade III) Tyranno SA fiber are listed in Table I. The thickness of the fiber coating was
0.8 m. To lower the sintering temperature and promote densification, -SiC nanopowder (Marketech International, Inc., Port Townsend, WA) was used for matrix formation, and Al2O3 (Sumitomo Chemical Industries, Ltd., Tokyo, Japan) and Y2O3 (Johnson Matthey, Inc., Cheshire, U.K.) were used as sintering aids. Typical compositions of the SiC nanopowder were 95% SiC, 1%–2% free carbon, and 1%–1.5% oxygen; the average particle size was 30 nm. Because the nanosized SiC powder used for matrix formation has a very high specific surface area (110 m2 /g), it was very difficult to modify the slurry to fit for the traditional filament￾winding method. In the present experiment, the PyC-coated fiber tows were first wound and fixed onto frames, to form aligned UD fiber sheets. These sheets were then impregnated by a polymer￾containing slurry (polycarbosilane (PCS) mixed with a filler, which was composed of SiC nanopowder and sintering additives) and pyrolyzed in an argon atmosphere for pretreatment. The polymer precursor of PCS (Nippon Carbon Co., Tokyo, Japan) was included to decrease the viscosity of the slurry and increase the wettability between the fibers and the matrix slurry. In this experiment, hexane was used as solvent. During polymer (with filler) impregnation and pyrolysis (PIP), the intrabundle matrix was compacted. The polymer-containing slurry needed to be carefully modified to increase the intrabundle infiltration effi￾ciency and to avoid matrix cracking. Because of the very high specific surface area (110 m2 /g) of nanopowder, higher SiC addition would greatly increase the viscosity of the slurry at a certain amount of solvent so that the infiltration efficiency was R. Naslain—contributing editor Manuscript No. 187345. Received November 7, 2001; approved October 10, 2002. This work was supported by CREST, Japan Science and Technology Corp. (JST). *Member, American Ceramic Society. † Present address: Shanghai Institute of Ceramics, Chinese Academy of Sciences, Shanghai 200050, People’s Republic of China. J. Am. Ceram. Soc., 86 [1] 26–32 (2003) 26 journal

anuary 200 Preparation of siC/siC Composites by Hot Pressing Table L. Properties of Tyranno SA Fiber(Grade ln C/Si atomic Diameter Elastic modulus SiC fiber Tyranno SA 1.08 7.5 3.10 1600 2.5 ~410 TData from Ube Industries, Ltd. shrinkag dui abundle matrix cracking the pyrolin gne farge previous report AL, O, and Y2 0, can be used as addities a decreased. If the filler content was too lo sintering monolithic SiC, via liquid-phase sintering. Under this addition in the present experiment was about 50 wt%. After PIP system, the temperature for liquid-phase formation between Al,o pretreatment, the sheets were cut into squares measuring "30 and Y,O3 can be lowered to 1700%-1800C by including a small mm X 30 mm, which were ready for further composite amount of SiO,, which inevitably exists as a surface layer on the tion. The cut squares were then coated with the matrix(nanopow Sic powders. However, densification is still not efficient when the der SiC with sintering additives) by dipping in the slurry, which temperature is 1750C, under 15 cation of monolithic SiC for the adopted matrix system. MPa pressure. At lower temperature(1720.C), the monolithic SiC 4 mm X 1.6 mm X 26 mm rectangles for three-point-bend testing probably because of the lower amount of liquid- phase formation (Model 5581, Instron Corp, Danvers, MA), with a crosshead At 1780C and 15 MPa, -99% of theoretical density was speed of 0.5 mm/min and a span of 18 mm. For comparison, tensile obtained, indicating that the presently selected material system an testing was also performed on some of the samples. The tensile- additive composition effectively promoted densification test specimens measured 3 mm X 1.6 mm X 30 mm On both ends For composite fabrication, the densification process slowed of the tensile bars, aluminum tabs were affixed to each side, using fter fiber inclusion in the matrix, as indicated in Fig. 1.A ee type of standard epoxy binder. The gauge length was designated relatively lower density was achieved under 15 MPa pressure at a be 15 mm. During the tensile test, cycle unloading-reloading given sintering temperature. Because effective densification could was applied to the tensile bars through the aluminum tabs on both not be obtained at lower temperature, even for monolithic SiC ends of the samples, which were clamped onto the fixture mounted fabrication, no composites were prepared at 1720.C and 15 MPa, on the test machine. The tensile strain was recorded using an although the density of such composites presumably would be extensometer fixed onto both sides of the gauge areas rather low under such conditions. To further enhance densification method. Both the polished cross sections and the fracture surfaces sintering temperature under the same pressure (either 15 or 20 microscopy(SEM). Further microstructural evaluation was con- Meanwhile, an increase in density at the same temperature under ducted by transmission electron microscopy (TEM) different pressures was also revealed, and this increase seemed more significant at a higher temperature, 1780.C, However, composite with a relative density of 90.8% was obtained at 1720C lIL. Results and 20 MPa, and a highly densified UD composite with a relative (I Densification Process density of 96.6% was achieved at 1780.C and 20 MP Before studying the fabrication of SiC/SiC composites, it is latrix itself i.e. without (2) Microstructural Evolution Figures 2 and 3 show the microstructures of polished cre sections for composites densified at 1720C and 20 MPa, and 1780.C and 20 MPa, respectively. Typically, no interbundle pores were observed, even in the composite sintered at 1720.C, under 20 MPa, as indicated in Figs. 2(a) and(b). Under the present preparation method, matrix-rich layers were clearly identified imilarly to the case for the densification of monolithic SiC, as 1780C, under 20 MPa, as shown in Fig 3 The magnified SEM micrographs in Figs. 2(b) and 3(a) show -M15 that micropores existed in the intrabundle areas for all of the C-15 composites, although both the number and size of pores decreased c-20 as sintering temperature and pressure increased. Even at 1780C and 20 MPa, many small pores were still distributed in the 700 1800 For the well-formed intrabundle matrix. little difference could Temperature C) be identified with variations of the sintering conditions. In the composite sintered at 1720.C, under 20 MPa, some consolidated Fig. 1. Effects of sintering conditions on the densification of the parts were loosely formed(shown by arrow in Fig. 2(c). On the monolithic sic and sic/SiC monolithic Sic hot. ther hand, at 1780.C, under 20 MPa, these types of intrabundle ressed under 15 MPa,(C-15)composites hot-pressed under 15 MPa, and regions were well consolidated, and fibers and matrix were tightly (C-20)composites hot-pressed under 20 MPa. bonded together (indicated by arrows in Fig. 3(b))

decreased. If the filler content was too low, the pyrolyzing process would cause intrabundle matrix cracking because of the large shrinkage during PCS to ceramic transformation. Suitable filler addition in the present experiment was about 50 wt%. After PIP pretreatment, the sheets were cut into squares measuring
30 mm  30 mm, which were ready for further composite prepara￾tion. The cut squares were then coated with the matrix (nanopow￾der SiC with sintering additives) by dipping in the slurry, which was mixed by attrition ball milling for 3 h. This process allowed a certain amount of interbundle matrix to form, and the fiber volume fraction could be controlled at
30 vol%, depending on the densification process used for the composites. After the PIP￾pretreated and matrix-coated squares had been dried, they were stacked in a graphite die and hot-pressed. The temperature was varied from 1720° to 1780°C, while the pressure was varied from 15 to 20 MPa. The holding time at the highest temperature was
1 h. Selection of the sintering conditions was based on the densifi￾cation of monolithic SiC for the adopted matrix system. The hot-pressed samples were subsequently cut and ground into 4 mm  1.6 mm  26 mm rectangles for three-point-bend testing (Model 5581, Instron Corp., Danvers, MA), with a crosshead speed of 0.5 mm/min and a span of 18 mm. For comparison, tensile testing was also performed on some of the samples. The tensile￾test specimens measured 3 mm  1.6 mm  30 mm. On both ends of the tensile bars, aluminum tabs were affixed to each side, using a type of standard epoxy binder. The gauge length was designated to be 15 mm. During the tensile test, cycle unloading–reloading was applied to the tensile bars through the aluminum tabs on both ends of the samples, which were clamped onto the fixture mounted on the test machine. The tensile strain was recorded using an extensometer fixed onto both sides of the gauge areas. The density of each sample was measured by the Archimedes method. Both the polished cross sections and the fracture surfaces after the bend and tensile tests were observed by scanning electron microscopy (SEM). Further microstructural evaluation was con￾ducted by transmission electron microscopy (TEM). III. Results (1) Densification Process Before studying the fabrication of SiC/SiC composites, it is necessary to study the sintering of the matrix itself, i.e., without fibers, to work out the proper HP conditions. According to a previous report,9 Al2O3 and Y2O3 can be used as additives for sintering monolithic SiC, via liquid-phase sintering. Under this system, the temperature for liquid-phase formation between Al2O3 and Y2O3 can be lowered to 1700°–1800°C by including a small amount of SiO2, which inevitably exists as a surface layer on the SiC powders. However, densification is still not efficient when the temperature is 1800°C, for micrometer-sized SiC powders. To avoid potential fiber degradation caused by temperature (especial￾ly under pressure), the selected temperature and pressure should be as low as possible. One possible way to lower the sintering temperature is to select highly active SiC nanopowder and care￾fully modify the composition of the sintering additives. Figure 1 shows variations of the density of the monolithic SiC used for matrix formation in the composites and the density of the composites under different temperatures and pressures. For mono￾lithic SiC, a higher density was obtained at 1750°C, under 15 MPa pressure. At lower temperature (1720°C), the monolithic SiC was still difficult to highly densify, even using a nanopowder, probably because of the lower amount of liquid-phase formation. At 1780°C and 15 MPa,
99% of theoretical density was obtained, indicating that the presently selected material system and additive composition effectively promoted densification. For composite fabrication, the densification process slowed after fiber inclusion in the matrix, as indicated in Fig. 1. A relatively lower density was achieved under 15 MPa pressure at a given sintering temperature. Because effective densification could not be obtained at lower temperature, even for monolithic SiC fabrication, no composites were prepared at 1720°C and 15 MPa, although the density of such composites presumably would be rather low under such conditions. To further enhance densification, a higher pressure of 20 MPa was adopted. With an increase in sintering temperature under the same pressure (either 15 or 20 MPa), the density of the composites increased continuously. Meanwhile, an increase in density at the same temperature under different pressures was also revealed, and this increase seemed more significant at a higher temperature, 1780°C. However, a composite with a relative density of 90.8% was obtained at 1720°C and 20 MPa, and a highly densified UD composite with a relative density of 96.6% was achieved at 1780°C and 20 MPa. (2) Microstructural Evolution Figures 2 and 3 show the microstructures of polished cross sections for composites densified at 1720°C and 20 MPa, and 1780°C and 20 MPa, respectively. Typically, no interbundle pores were observed, even in the composite sintered at 1720°C, under 20 MPa, as indicated in Figs. 2(a) and (b). Under the present preparation method, matrix-rich layers were clearly identified. Similarly to the case for the densification of monolithic SiC, as indicated in Fig. 1, these layers were well sintered, especially at 1780°C, under 20 MPa, as shown in Fig. 3(c). The magnified SEM micrographs in Figs. 2(b) and 3(a) show that micropores existed in the intrabundle areas for all of the composites, although both the number and size of pores decreased as sintering temperature and pressure increased. Even at 1780°C and 20 MPa, many small pores were still distributed in the intrabundle regions. For the well-formed intrabundle matrix, little difference could be identified with variations of the sintering conditions. In the composite sintered at 1720°C, under 20 MPa, some consolidated parts were loosely formed (shown by arrow in Fig. 2(c)). On the other hand, at 1780°C, under 20 MPa, these types of intrabundle regions were well consolidated, and fibers and matrix were tightly bonded together (indicated by arrows in Fig. 3(b)). Table I. Properties of Tyranno SA Fiber (Grade III)† SiC fiber C/Si atomic ratio Diameter (m) Density (g/cm3 ) Filaments/yarn Tensile strength (GPa) Elastic modulus (GPa) Tyranno SA 1.08 7.5 3.10 1600
2.5
410 † Data from Ube Industries, Ltd. Fig. 1. Effects of sintering conditions on the densification of the monolithic SiC and SiC/SiC composites: (M-15) monolithic SiC hot￾pressed under 15 MPa, (C-15) composites hot-pressed under 15 MPa, and (C-20) composites hot-pressed under 20 MPa. January 2003 Preparation of SiC/SiC Composites by Hot Pressing 27

Journal of the American Ceramic Sociery-Dong et al. Vol. 86. No. I Matrix rich lay Oum (b) (b) Interbundle matrix PyC interphase Intrabundle matrIx nOun (c) Intrabundle matrix The HP-sintered SiC interlayer matrix m Fig. 2. SEM micrographs of the polished cross section for the composite Fig 3. SEM micrographs of the polished cross section for the composite sintered at 1780C. 20 MPa 3) Physical and Mechanical Properties Meanwhile, the values for strength, PLS, and elastic modulus Some physical and mechanical properties of the composites nder different sintering conditions, as obtained from the bend test 78 dually increased, and the highest values were obtained at 80C and 20 MPa. The bending strength under those sintering are listed in Table Il. The proportional limit stress(PLS) and the conditions was >700 MPa ultimate bending strength (UBS)were determined from the load- Typical stress-displacement curves derived from the bend test displacement curves according to ASTM C 1341-97. The Pls are shown in Fig. 4. These curves demonstrate noncatastrophic was the stress corresponding to a 0.01% offset strain. As shown in fracture behavior. The curves have a similar appearance, with little Table Il, when the sintering temperature and pressure increase difference in characteristics such as the uBs. the Pls. and the the porosity of the composites decreased simultaneously. At elastic modulus, indicating consistency with the data listed in 1780.C and 20 MPa, the porosity reached a lowest value of 3.4% Table ll

(3) Physical and Mechanical Properties Some physical and mechanical properties of the composites under different sintering conditions, as obtained from the bend test, are listed in Table II. The proportional limit stress (PLS) and the ultimate bending strength (UBS) were determined from the load– displacement curves according to ASTM C 1341-97.10 The PLS was the stress corresponding to a 0.01% offset strain. As shown in Table II, when the sintering temperature and pressure increased, the porosity of the composites decreased simultaneously. At 1780°C and 20 MPa, the porosity reached a lowest value of 3.4%. Meanwhile, the values for strength, PLS, and elastic modulus gradually increased, and the highest values were obtained at 1780°C and 20 MPa. The bending strength under those sintering conditions was 700 MPa. Typical stress–displacement curves derived from the bend test are shown in Fig. 4. These curves demonstrate noncatastrophic fracture behavior. The curves have a similar appearance, with little difference in characteristics such as the UBS, the PLS, and the elastic modulus, indicating consistency with the data listed in Table II. Fig. 2. SEM micrographs of the polished cross section for the composite sintered at 1720°C, 20 MPa. Fig. 3. SEM micrographs of the polished cross section for the composite sintered at 1780°C, 20 MPa. 28 Journal of the American Ceramic Society—Dong et al. Vol. 86, No. 1

anuary 2003 Preparation of siC/siC Composites by Hot Pressing Table IL. Effects of Experimental Conditions on the Physical and Mechanical Properties of the Composites 1720°C 1750° 1780°C essure(MPa) 2.77±0 2.79±0.032.83±0.012.85±0.032.93±0. porosity(%) Bending strength(MPa) 560±10 630±36 650±14 680±20 710±17 Proportional limit stress(MPa) 350+27 520±21 lastic modulus(GPa 160±7 180±13 180±7 190±8 220±8 Tensile tests also were conducted on some of the composites, increases, from 220 to 280 GPa. At the same time, the stress-strain and typical mechanical properties of ultimate tensile strength and curves also demonstrate typical behavior, in that each curv corresponding PLs for those composites are indicated in Fig. 5 potentially reflects elastic deformation, matrix cracking, elasti Data analysis indicated a similar variation trend for both the tensile elongation of the fibers, and individual fiber breaks, as described test and the bend test, with increases in sintering temperature from in the literature. However, all of these stress-strain curves have 1720 to 1780%C. under 20 MPa. Both the ultimate tensile strength a wide curved domain. Figure 7 reveals the variation of the elastic and the Pls were highest at 1780.C and 20 MPa. Under these modulus(normalized by initial modulus of each composite)during sintering conditions, the ultimate tensile strength was 400 MP and the pls was 210 MPa Stress-strain curves from the tensile test with unloading- reloading cycles are shown in Fig. 6. Comparison of these curves shows that an increase in temperature narrows the width of th hysteresis loops and decreases the permanent strain at zero load, a 350 or samples under identical pressures of 20 MPa. These features 3 rrespond to the difference in fiber/matrix bond lation of the elastic modulus from the slope of the initial linear 9 250 ortion of the reloading curve shown in Fig. 6 also indicates a characteristic unique to each composite: As the temperature 150 increases from 1720 to 1780 C. the tensile elastic modulus also 800 1780℃/20MPa Tensile strain(%) 1720C/20MPa m至+巴5 150 Displacement(mm) Fig. 4. Typical stress-displacement curves from bend test for the 0 04 composites sintered at different conditions. Tensile strain(% (c) 250 Ultimate tensile stress Proportional limit stress 1700 1750 0.1 0.2 0.3 0.4 Te Temperature) Fig. 6. Tensile in curves with unloading-reloading cycles for 5. Typical mechanical rties of the composites sintered at opposites sintered at different conditions: (a)1720C/20 MPa,(b) rent temperatures under an identical pressure of 20 MPa. 1750°c/20MPa,and(c)1780°C/20MPa

Tensile tests also were conducted on some of the composites, and typical mechanical properties of ultimate tensile strength and corresponding PLS for those composites are indicated in Fig. 5. Data analysis indicated a similar variation trend for both the tensile test and the bend test, with increases in sintering temperature from 1720° to 1780°C, under 20 MPa. Both the ultimate tensile strength and the PLS were highest at 1780°C and 20 MPa. Under these sintering conditions, the ultimate tensile strength was
400 MPa, and the PLS was
210 MPa. Stress–strain curves from the tensile test with unloading– reloading cycles are shown in Fig. 6. Comparison of these curves shows that an increase in temperature narrows the width of the hysteresis loops and decreases the permanent strain at zero load, for samples under identical pressures of 20 MPa. These features correspond to the difference in fiber/matrix bonding.11,12 Calcu￾lation of the elastic modulus from the slope of the initial linear portion of the reloading curve shown in Fig. 6 also indicates a characteristic unique to each composite: As the temperature increases from 1720° to 1780°C, the tensile elastic modulus also increases, from 220 to 280 GPa. At the same time, the stress–strain curves also demonstrate typical behavior, in that each curve potentially reflects elastic deformation, matrix cracking, elastic elongation of the fibers, and individual fiber breaks, as described in the literature.11 However, all of these stress–strain curves have a wide curved domain. Figure 7 reveals the variation of the elastic modulus (normalized by initial modulus of each composite) during Table II. Effects of Experimental Conditions on the Physical and Mechanical Properties of the Composites 1720°C 1750°C 1780°C Pressure (MPa) 20 15 20 15 20 Density (g/cm3 ) 2.77  0.02 2.79  0.03 2.83  0.01 2.85  0.03 2.93  0.02 Porosity (%) 9.2 7.8 6.6 5.7 3.4 Bending strength (MPa) 560  10 630  36 650  14 680  20 710  17 Proportional limit stress (MPa) 350  27 430  22 440  30 450  15 520  21 Elastic modulus (GPa) 160  7 180  13 180  7 190  8 220  8 Fig. 4. Typical stress–displacement curves from bend test for the composites sintered at different conditions. Fig. 5. Typical mechanical properties of the composites sintered at different temperatures under an identical pressure of 20 MPa. Fig. 6. Tensile stress–strain curves with unloading–reloading cycles for the composites sintered at different conditions: (a) 1720°C/20 MPa, (b) 1750°C/20 MPa, and (c) 1780°C/20 MPa. January 2003 Preparation of SiC/SiC Composites by Hot Pressing 29

Journal of the American Ceramic Sociery-Dong et al. Vol. 86. No. I (a) 0.6 EVAE 02 0005010.150202503035 Strain(%) Fig. 7. Normalized tensile modulus versus tensile strain for the composites sintered at different conditions: (a)1720.C/20 MPa, (b)1750C/20 MPa, ar (c)1780.C/20 MPa; (E) elastic modulus given by m tangent modulus; (Eo) the initial elastic modulus of each composite, (Ed the fiber Youngs modulus; (o the fiber volume fraction unloading-reloading cycles, using a similar method described in so that fiber rearrangement is continuously enhanced under the literature. 2 As shown, all of the elastic moduli of the In the present case, a highly densified composite was composites decrease gradually, which suggests the accumulation at 1780.C, under 20 MPa pressure(as indicated in Fig. 1) of matrix cracks during the tensile cycling. The moduli at ultimate However, during this process, higher external pressure simulta- failure of the composites are still over the value of ers indicating neously exerting force on the surfaces of the fibers could cause that the debonding between fiber and matrix has not fully accom- potential damage to the fibers with an increase of both temperature plished and the saturation of matrix cracking has not occurred. and pressure Because densification and microstructural evolution are different As discussed earlier, intrabundle-matrix formation is highly under different sintering conditions, the characteristics of the dependent on impregnation efficiency during PIP pretreatment. stress-strain curves also demonstrate their inherent features Because the nanopowder has a very high specific surface area, infiltration of the matrix to the fiber bundles becomes a challenge (4 Fractography for this process. Higher infiltration efficiency can be obtained b Figure 8 shows the fracture surfaces of composites prepared at modifying the ratio of polymer(PCS) to filler(nanopowder SiC different sintering temperatures, from 1720 to 1780.C, under 20 with sintering additives) and the solvent content. However, for a MPa pressure. As temperature increases, the fracture patterns, polymer precursor, the conversion of the PCs to an amorphous mainly identified by fiber pullout, clearly change. At lower pyrolyzed product, followed by crystallization at high temperature, sintering temperature(1720.C), long fiber pullout dominates the is accompanied by a large volume contraction. As a result, pores racture surface, as shown in Fig 8(a). In this particular composite, occur in the intrabundle areas. During sintering(especially under the carbon coating has attached mainly onto the fiber surface pressure), some of the pores may be removed by fiber rearrange- luring fracture, indicating that matrix cracking has propagated ment, but some remain. along the matrix/carbon interface. As the temperature increases tions of the composites sintered at different temperatures, from sintering temperature(1780%C), short fiber pullout dominates the 720 to 1780C, under 20 MPa pressure, reveals that apparent acture surface, as shown in Fig. 8(c) differences are distinguishable only in the intrabundle area. The In many cases, debond occurs at the fiber/carbon interface, as pore size and the number of pores decrease when the temperature indicated in Fig. 9. Even in this composite, relatively long fiber increases. At the same time, as shown by the magnified micro- pullout also has been identified, especially in the inner fiber layers, graphs in Figs. 2(c)and 3(b), the intrabundle matrix densifies even dicating that some weak fiber/matrix bonding still exists, allow- better as the temperature increases. Little difference is identifiable ing long crack deflection. between the interbundle matrix and the intrabundle areas. the densification of monolithic SiC, as shown in Fig. 1, implies that the matrix is well densified at 1750.C, under 15 MPa pressure Further densification of the composites is attributed mainly to fiber Efect of sintering Conditions on the rearrangement and intrabundle-matrix (containing polymer Densification Process derived products)densification, and, in this case, higher tempera- ture(up to 1780C) and higher pressure (up to 20 MPa)more The densification process for the composites can be divided into effectively promote densification of the composite matrix densification. Densification of the interbundle matrix is similar to that for monolithic Sic and is a temperature-dependent (2) Efect of Sintering Conditions on Mechanical Behavior process, as shown in Fig. 1. At>1750C, under 15 MPa pressure, Because the densification and microstructural evolution of the a highly densified interbundle matrix should be achieved, even composites change with variations in sintering conditions, inter fiber inclusion slows the densification process action between the fibers and the matrix, and thus the mechanical Intrabundle-matrix formation is highly controlled by impregnation behavior of the composites are directly affected. As discussed during PIP pretreatment. Because the PIP-formed products have arlier, composites sintered under lower temperature and/or pres- some residual carbon, resulting from the inclusion of the PCs sure have lower density and higher porosity, and the micropores lymer precursor, the sinterability is more or less inhibited. At are concentrated mainly in the intrabundle areas. These types of the same time. fiber re ement is also dependent on the eatures benefit crack initiation during bend or tensile testing, so sinterability of the intrabundle matrix. At higher temperature, that the composites exhibit relatively lower PLS. Meanwhile, liquid-phase formation more effectively promotes grain-boundary because either low temperature or low pressure leads to the

unloading–reloading cycles, using a similar method described in the literature.12 As shown, all of the elastic moduli of the composites decrease gradually, which suggests the accumulation of matrix cracks during the tensile cycling. The moduli at ultimate failure of the composites are still over the value of Ef Vf , indicating that the debonding between fiber and matrix has not fully accom￾plished and the saturation of matrix cracking has not occurred. Because densification and microstructural evolution are different under different sintering conditions, the characteristics of the stress–strain curves also demonstrate their inherent features. (4) Fractography Figure 8 shows the fracture surfaces of composites prepared at different sintering temperatures, from 1720° to 1780°C, under 20 MPa pressure. As temperature increases, the fracture patterns, mainly identified by fiber pullout, clearly change. At lower sintering temperature (1720°C), long fiber pullout dominates the fracture surface, as shown in Fig. 8(a). In this particular composite, the carbon coating has attached mainly onto the fiber surface during fracture, indicating that matrix cracking has propagated along the matrix/carbon interface. As the temperature increases under the same pressure, the pulled-out fibers shorten. At higher sintering temperature (1780°C), short fiber pullout dominates the fracture surface, as shown in Fig. 8(c). In many cases, debond occurs at the fiber/carbon interface, as indicated in Fig. 9. Even in this composite, relatively long fiber pullout also has been identified, especially in the inner fiber layers, indicating that some weak fiber/matrix bonding still exists, allow￾ing long crack deflection. IV. Discussion (1) Effect of Sintering Conditions on the Densification Process The densification process for the composites can be divided into two main parts: interbundle-matrix densification and intrabundle￾matrix densification. Densification of the interbundle matrix is similar to that for monolithic SiC and is a temperature-dependent process, as shown in Fig. 1. At 1750°C, under 15 MPa pressure, a highly densified interbundle matrix should be achieved, even though fiber inclusion slows the densification process. Intrabundle-matrix formation is highly controlled by impregnation during PIP pretreatment. Because the PIP-formed products have some residual carbon, resulting from the inclusion of the PCS polymer precursor,13 the sinterability is more or less inhibited. At the same time, fiber rearrangement is also dependent on the sinterability of the intrabundle matrix. At higher temperature, liquid-phase formation more effectively promotes grain-boundary sliding, so that fiber rearrangement is continuously enhanced under pressure. In the present case, a highly densified composite was obtained at 1780°C, under 20 MPa pressure (as indicated in Fig. 1). However, during this process, higher external pressure simulta￾neously exerting force on the surfaces of the fibers could cause potential damage to the fibers with an increase of both temperature and pressure. As discussed earlier, intrabundle-matrix formation is highly dependent on impregnation efficiency during PIP pretreatment. Because the nanopowder has a very high specific surface area, infiltration of the matrix to the fiber bundles becomes a challenge for this process. Higher infiltration efficiency can be obtained by modifying the ratio of polymer (PCS) to filler (nanopowder SiC with sintering additives) and the solvent content. However, for a polymer precursor, the conversion of the PCS to an amorphous pyrolyzed product, followed by crystallization at high temperature, is accompanied by a large volume contraction. As a result, pores occur in the intrabundle areas. During sintering (especially under pressure), some of the pores may be removed by fiber rearrange￾ment, but some remain. Comparison of the microstructures of the polished cross sec￾tions of the composites sintered at different temperatures, from 1720° to 1780°C, under 20 MPa pressure, reveals that apparent differences are distinguishable only in the intrabundle area. The pore size and the number of pores decrease when the temperature increases. At the same time, as shown by the magnified micro￾graphs in Figs. 2(c) and 3(b), the intrabundle matrix densifies even better as the temperature increases. Little difference is identifiable between the interbundle matrix and the intrabundle areas. The densification of monolithic SiC, as shown in Fig. 1, implies that the matrix is well densified at 1750°C, under 15 MPa pressure. Further densification of the composites is attributed mainly to fiber rearrangement and intrabundle-matrix (containing polymer￾derived products) densification, and, in this case, higher tempera￾ture (up to 1780°C) and higher pressure (up to 20 MPa) more effectively promote densification of the composite. (2) Effect of Sintering Conditions on Mechanical Behavior Because the densification and microstructural evolution of the composites change with variations in sintering conditions, inter￾action between the fibers and the matrix, and, thus, the mechanical behavior of the composites are directly affected. As discussed earlier, composites sintered under lower temperature and/or pres￾sure have lower density and higher porosity, and the micropores are concentrated mainly in the intrabundle areas. These types of features benefit crack initiation during bend or tensile testing, so that the composites exhibit relatively lower PLS. Meanwhile, because either low temperature or low pressure leads to the Fig. 7. Normalized tensile modulus versus tensile strain for the composites sintered at different conditions: (a) 1720°C/20 MPa, (b) 1750°C/20 MPa, and (c) 1780°C/20 MPa; (E) elastic modulus given by minimum tangent modulus; (E0) the initial elastic modulus of each composite; (Ef ) the fiber Young’s modulus; (Vf ) the fiber volume fraction. 30 Journal of the American Ceramic Society—Dong et al. Vol. 86, No. 1

anuary 200 Preparation of siC/siC Composites by Hot Pressing Fig. 9. Fracture surface of the composite sintered at 1780C under 20 MPa showing the typical characteristics for short fiber pull-out. PR both bend and tensile tests. Relatively lower strength is obtained for composites sintered at 1720C, under 20 MPa pressure, as shown in Table II and Fig. 5 Composites sintered under higher temperature and pressure (e.g, 1780C and 20 MPa) contain both a highly densified interlayer matrix and a well-formed intrabundle matrix with lowe portion of the loading curves shown in Fig. 4 for the bend test and in Fig. 6(c)for the tensile test. Compared with composites sintered at 1720 C and 20 MPa, the well-densified matrixes and enhanced 500um fiber/matrix interfaces of composites sintered at higher tempera ure(1780%C), under identical pressures, can withstand a higher stress for elastic deformation this behavior also contributes to a high ability of the composites to arrest crack initiation, so that the composites exhibit a higher PLs. At the same time, the rather narrow hysteresis loops and decreased perm at load also suggest stronger fiber/matrix bonding. This stronger bonding allows short crack deflection, leading to short fibe pullout(Figs. 8(c)and 9), and simultaneously aids load transfer 停多 Tyranno SA 710±68nm (a)1720°c/20MPa,(b)1750°c/20MPa,and(c)1780°0MPa tion of a less-densified pecially for the intrabundle matrix. a relatively weak fiber/matrix interface forms. This inter- ace aids in debonding and frictional sliding under load 4-l Characterization of the unloading-rel curve as illustrated in 你念 Fig. 6(a), reveals a typical feature, which consists of a ermanent strain at zero load. As commonly recognized SiC/SiC composites 17, 18 weak fiber/matrix bonding allows long rack deflection along the interface, so that long fiber pullout occurs ever eak fiber/matrix bonding is beneficial for toughening composites, it is simultaneously detri ental for strengthening Consistent results have been revealed by Fig 10. TEM observation of the composite sintered at 1780.C, 20 MPa

formation of a less-densified matrix, especially for the intrabundle matrix, a relatively weak fiber/matrix interface forms. This inter￾face aids in debonding and frictional sliding under load.14–16 Characterization of the unloading–reloading curve, as illustrated in Fig. 6(a), reveals a typical feature, which consists of a large permanent strain at zero load. As commonly recognized for SiC/SiC composites,17,18 weak fiber/matrix bonding allows long crack deflection along the interface, so that long fiber pullout occurs. However, even though weak fiber/matrix bonding is beneficial for toughening composites, it is simultaneously detri￾mental for strengthening. Consistent results have been revealed by both bend and tensile tests. Relatively lower strength is obtained for composites sintered at 1720°C, under 20 MPa pressure, as shown in Table II and Fig. 5. Composites sintered under higher temperature and pressure (e.g., 1780°C and 20 MPa) contain both a highly densified interlayer matrix and a well-formed intrabundle matrix with lower porosity. This microstructural evolution contributes a higher elas￾tic modulus to the composite, as indicated by the slope of the linear portion of the loading curves shown in Fig. 4 for the bend test and in Fig. 6(c) for the tensile test. Compared with composites sintered at 1720°C and 20 MPa, the well-densified matrixes and enhanced fiber/matrix interfaces of composites sintered at higher tempera￾ture (1780°C), under identical pressures, can withstand a higher stress for elastic deformation. This behavior also contributes to a high ability of the composites to arrest crack initiation, so that the composites exhibit a higher PLS. At the same time, the rather narrow hysteresis loops and decreased permanent strain at zero load also suggest stronger fiber/matrix bonding. This stronger bonding allows short crack deflection, leading to short fiber pullout (Figs. 8(c) and 9), and simultaneously aids load transfer Fig. 8. Fractography of the composites at different sintering conditions: (a) 1720°C/20 MPa, (b) 1750°C/20 MPa, and (c) 1780°C/20 MPa. Fig. 9. Fracture surface of the composite sintered at 1780°C under 20 MPa showing the typical characteristics for short fiber pull-out. Fig. 10. TEM observation of the composite sintered at 1780°C, 20 MPa. January 2003 Preparation of SiC/SiC Composites by Hot Pressing 31

Journal of the American Ceramic Sociery-Dong et al. Vol. 86. No. I from the matrix to the fibers. Combined with higher densification relatively strong trix bonding. This strong bonding, of the matrix, higher strength is obtained (Table Il and Fig. 5) combined with the med matrix (in both the intrabundle As discussed in the literature. 4, 6, 9-2 fiber. matrix. and and the interbund provided the composite with interface are the important factors for determining the properties of proved mechanical propert opposites. In the present experiment, potential reactions between SiO, and carbon(either from the polymer-derived matrix or the carbon coating), and between Al,O, and Y2O3, could change the References omposition of the matrix during high-temperature sintering. Although carbon interphase still maintains well from the TEM K Nakano, A. Kamiya, H Ogawa, and Y. Nishino, "Fabrication and Mechanical bservation, as shown in Fig. 10, this change in composition would Properties of Carbon Fiber Reinforced Silicon Carbide Composites, "J. Ceran. Soc. Jpm,10041472-75(1992) components of the composite(fiber, matrix, and interphase). Such an effect is difficult to estimate; in the present case, further L. Porte and A. Sartre, "Evidence for a Silicon Oxycarbide Phase in the Nicalon Silicon Carbide Fiber, J. Mater Sci., 24, 271-75(1989) haracterization is necessary However, for the composites sintered under higher temperature J Dixmier, J. L. Miquel, H. Hammel, nd pressure(1780C and 20 MPa), individual fiber deformation is by EXAFS S P. Legran also detectable in a few cases, indicating that potential fiber creep Some Additional Methods, J. Mater. Sci., 24, 1503-12(1989)- Bodet, N. Jia, and R. E. Tressler, ""Microstructural Instability and the Resultant deformation might have occurred. For the present experiment, strength of Si-C-O(Nicalon) and Si-N-C-O(HPZ) Fibres, J. Eur. Ceram. Soc. creep deformation does not seem to have significantly affected the 6653-64(1996 performances of the composites. Generally, the fibers were well protected by the carbon coating and contributed high performance (95rbide-Based Fibre with a Low Oxygen Content,".Mater. Sci, 30, 661-77 and noncatastrophic fracture behavior to the composites. Improv 7T. Ishikawa, Y. Kohtoku, K. Kumagawa, T. Yamamura, and T. Nagasawa, ing intrabundle-matrix formation and increasing the fiber volume "High-Strength Alkali-Resistant Sintered SiC Fibre Stable to 2200oC, "Nature fraction might further enhance the composite characteristics, and London),391,773-75(1998) these goals will become the focus of continuing efforts. AS. M. Dong, G. Cholon, C. Labrugere, M. Lahaye, A. Guette, R. Naslain, an D. L. Jiang, "Characterization of Some Advanced Sic-Based Ceramic Fibers, J. 36,2371-8102001 M. A. Mulla and V. D. Krstic, Low-Temperature Pressureless V. Conclusions m Oxide and Yttrium Oxide Additions. Am. Ceram. Soc.Bm,713]439-43(1991) Unidirectional SiC/SiC composites were successfully fabricated 10-Standard Test Method for Flexural Properties of Continuous Fiber-Reinforced by HP, via liquid-phase sintering, at temperatures from 1720%to Advanced Ceramic Composites, ASTM C 1341-97, Pp. 509-26. American Society 1780C, under pressures from 15 to 20 MPa, using advanced for Testing and Materials, West Conshohocken, P/ S.Bertrand, P Forio, R. Pailler, and J. Lamon, "Hi-Nicalon/SiC Minicomposites a well-crystallized microstructure. Carbon coatings protected the 99)2465-73(1999). with(Pyrocarbon/SiC)e Nanoscale Multilayered Interphases, "J.Am. Ceram Soc., 82 fibers well from damage under harsh sintering conditions and Bertrand, R. Pailler, and J. Lamon, "Influence of Strong Fiber/c ontributed to high performance in the composites Mechanical Behavior and Lifetime of Hi-Nicalon/( PyC/SiC),SiC Minicomposites, The densification process as highly dependent on sintering Sor Soc,844]787-94(2001 tural Evolutions from e mperature and pressure in the present investigation. As temper- Polycarbosilane to SiC Ceramic,J, Mater. Sci., 25, 3886-93(1990). sarath, "The Role ultimate bending or tensile strength, PLs, and elastic module Interface in Ceramic Composites, "Am Ceram. Soc. Bull, 68 [2] improved. At 1780C, under 20 MPa pressure, a highly densified P. D. Jero, R.J. Kerans, and T A. Parthasarathy, "Effect of Interfacial Roughness omposite, with a density of 2.93 g/cm and a porosity of 3. 4%, on the Frictional Stress Measured Using Pushout Tests,"JAm Ceram Soc, 74(11] as obtained. The bending strength of this composite was >700 2795-801(I99)) APa, and the tensile strength was 400 MPa. Even though high mon, R. Naslain, E. L. Curzio, M. K. Ferber, and T. M mperature and high pressure can cause fiber damage and lead to Single-Fiber Push-Out Tests,"J. Am. Ceram Soc, 81 14)965-78(1998 degradation, strengthening and toughening by fiber reinforcement 7F. Rebillat, J. Lamon, R. Naslain, E. L. Curzio, M. K. Ferber, and T. M till predominated for the present experiment. All of the compos- Besmann,"Properties of Multilayered Interphases in SiC/SiC Chemical-Vapor- tes prepared under the various conditions demonstrated non- Infiltrated Composites with""and"Strong'Interfaces, " J. Am Ceram Soc.,81 catastrophic fracture behavior 912315-26(1998 F. Rebillat, J. Lamon, and A. Guette At lower temperature and/or pressure, long fiber pullout was the to SiC/Sic Composites with a BN Interp nant fracture behavior. Characterization of the hysteresi ni,"Methodology for Relating the ps derived from the stress-strain curves of the tensile test with opposites to Constituent Prop- unloading-reloading cycles confirmed that a relatively weak fiber/ erties,J.Am. Ceram Soc., 77 [6] 1425-35(1994 M matrix interface had formed. These composites exhibited relatively ments and the Constituent Properties of Ceramic Matrix( low strength. On the other hand, at high temperature and/or pressure, short fiber pullout became the main fracture behavior 吗s. 7.-9.3.32.p9a The narrow hysteresis loops and reduced permanent strain at Constituent Properties of Ceramic Matri tes: Il, Experimental Studies on Unidirectional Materials, J. Am. Ceram Soc., 78 zero load. as derived from the stress-strain curves. suggest 012721-31(1995)

from the matrix to the fibers. Combined with higher densification of the matrix, higher strength is obtained (Table II and Fig. 5). As discussed in the literature,14,16,19–21 fiber, matrix, and interface are the important factors for determining the properties of composites. In the present experiment, potential reactions between SiO2 and carbon (either from the polymer-derived matrix or the carbon coating), and between Al2O3 and Y2O3, could change the composition of the matrix during high-temperature sintering. Although carbon interphase still maintains well from the TEM observation, as shown in Fig. 10, this change in composition would affect the fiber (interphase)/matrix interaction and the thermal residual stress related to the coefficient of thermal expansion of the components of the composite (fiber, matrix, and interphase). Such an effect is difficult to estimate; in the present case, further characterization is necessary. However, for the composites sintered under higher temperature and pressure (1780°C and 20 MPa), individual fiber deformation is also detectable in a few cases, indicating that potential fiber creep deformation might have occurred. For the present experiment, creep deformation does not seem to have significantly affected the performances of the composites. Generally, the fibers were well protected by the carbon coating and contributed high performance and noncatastrophic fracture behavior to the composites. Improv￾ing intrabundle-matrix formation and increasing the fiber volume fraction might further enhance the composite characteristics, and these goals will become the focus of continuing efforts. V. Conclusions Unidirectional SiC/SiC composites were successfully fabricated by HP, via liquid-phase sintering, at temperatures from 1720° to 1780°C, under pressures from 15 to 20 MPa, using advanced Tyranno SA SiC fiber with a near-stoichiometric composition and a well-crystallized microstructure. Carbon coatings protected the fibers well from damage under harsh sintering conditions and contributed to high performance in the composites. The densification process was highly dependent on sintering temperature and pressure in the present investigation. As temper￾ature and pressure increased, all of the mechanical properties, i.e., ultimate bending or tensile strength, PLS, and elastic modulus, improved. At 1780°C, under 20 MPa pressure, a highly densified composite, with a density of 2.93 g/cm3 and a porosity of 3.4%, was obtained. The bending strength of this composite was 700 MPa, and the tensile strength was
400 MPa. Even though high temperature and high pressure can cause fiber damage and lead to degradation, strengthening and toughening by fiber reinforcement still predominated for the present experiment. All of the compos￾ites prepared under the various conditions demonstrated non￾catastrophic fracture behavior. At lower temperature and/or pressure, long fiber pullout was the predominant fracture behavior. Characterization of the hysteresis loops derived from the stress–strain curves of the tensile test with unloading–reloading cycles confirmed that a relatively weak fiber/ matrix interface had formed. These composites exhibited relatively low strength. On the other hand, at high temperature and/or pressure, short fiber pullout became the main fracture behavior. The narrow hysteresis loops and reduced permanent strain at zero load, as derived from the stress–strain curves, suggest relatively strong fiber/matrix bonding. This strong bonding, combined with the well-formed matrix (in both the intrabundle and the interbundle areas), provided the composite with im￾proved mechanical properties. References 1 K. Nakano, A. Kamiya, H. Ogawa, and Y. Nishino, “Fabrication and Mechanical Properties of Carbon Fiber Reinforced Silicon Carbide Composites,” J. Ceram. Soc. Jpn., 100 [4] 472–75 (1992). 2 K. Park and T. Vasilos, “Processing, Microstructure and Mechanical Properties of Hot Pressed SiC Continuous Fibre/SiC Composites,” J. Mater. Sci., 32, 295–300 (1997). 3 L. Porte and A. Sartre, “Evidence for a Silicon Oxycarbide Phase in the Nicalon Silicon Carbide Fiber,” J. Mater Sci., 24, 271–75 (1989). 4 C. Laffon, A. M. Flank, P. Lagarde, M. Laridjani, R. Hagege, P. Olry, J. Cotteret, J. Dixmier, J. L. Miquel, H. Hammel, and A. P. Legrand, “Study of Nicalon-Based Ceramic Fibers and Powders by EXAFS Spectrometry, X-ray Diffractometry and Some Additional Methods,” J. Mater. Sci., 24, 1503–12 (1989). 5 R. Bodet, N. Jia, and R. E. Tressler, “Microstructural Instability and the Resultant Strength of Si-C-O (Nicalon) and Si-N-C-O (HPZ) Fibres,” J. Eur. Ceram. Soc., 16 [6] 653–64 (1996). 6 R. Bodet, X. Bourrat, J. Lamon, and R. Naslain, “Tensile Creep Behaviour of a Silicon Carbide-Based Fibre with a Low Oxygen Content,” J. Mater. Sci., 30, 661–77 (1995). 7 T. Ishikawa, Y. Kohtoku, K. Kumagawa, T. Yamamura, and T. Nagasawa, “High-Strength Alkali-Resistant Sintered SiC Fibre Stable to 2200°C,” Nature (London), 391, 773–75 (1998). 8 S. M. Dong, G. Chollon, C. Labrugere, M. Lahaye, A. Guette, R. Naslain, and D. L. Jiang, “Characterization of Some Advanced SiC-Based Ceramic Fibers,” J. Mater. Sci., 36, 2371–81 (2001). 9 M. A. Mulla and V. D. Krstic, “Low-Temperature Pressureless Sintering of -Silicon Carbide with Aluminum Oxide and Yttrium Oxide Additions,” Am. Ceram. Soc. Bull., 70 [3] 439–43 (1991). 10“Standard Test Method for Flexural Properties of Continuous Fiber-Reinforced Advanced Ceramic Composites,” ASTM C 1341-97, pp. 509–26. American Society for Testing and Materials, West Conshohocken, PA, 2000. 11S. Bertrand, P. Forio, R. Pailler, and J. Lamon, “Hi-Nicalon/SiC Minicomposites with (Pyrocarbon/SiC)n Nanoscale Multilayered Interphases,” J. Am. Ceram. Soc., 82 [9] 2465–73 (1999). 12S. Bertrand, R. Pailler, and J. Lamon, “Influence of Strong Fiber/Coating on the Mechanical Behavior and Lifetime of Hi-Nicalon/(PyC/SiC)n/SiC Minicomposites,” J. Am. Ceram. Soc., 84 [4] 787–94 (2001). 13G. D. Soraru, F. Babonneau, and J. D. Mackenzie, “Structural Evolutions from Polycarbosilane to SiC Ceramic,” J. Mater. Sci., 25, 3886–93 (1990). 14R. J. Kerans, R. S. Hay, N. J. Pagano, and T. A. Parthasarathy, “The Role of the Fiber-Matrix Interface in Ceramic Composites,” Am. Ceram. Soc. Bull., 68 [2] 429–42 (1989). 15P. D. Jero, R. J. Kerans, and T. A. Parthasarathy, “Effect of Interfacial Roughness on the Frictional Stress Measured Using Pushout Tests,” J. Am. Ceram. Soc., 74 [11] 2793–801 (1991). 16F. Rebillat, J. Lamon, R. Naslain, E. L. Curzio, M. K. Ferber, and T. M. Besmann, “Interfacial Bond Strength in SiC/SiC Composite Materials, As studied by Single-Fiber Push-Out Tests,” J. Am. Ceram. Soc., 81 [4] 965–78 (1998). 17F. Rebillat, J. Lamon, R. Naslain, E. L. Curzio, M. K. Ferber, and T. M. Besmann, “Properties of Multilayered Interphases in SiC/SiC Chemical-Vapor￾Infiltrated Composites with ‘Weak’ and ‘Strong’ Interfaces,” J. Am. Ceram. Soc., 81 [9] 2315–26 (1998). 18F. Rebillat, J. Lamon, and A. Guette, “The Concept of a Strong Interface Applied to SiC/SiC Composites with a BN Interphase,” Acta Mater., 48, 4609–18 (2000). 19A. G. Evans, J. M. Domergue, and E. Vagaggini, “Methodology for Relating the Tensile Constitutive Behavior of Ceramics-Matrix Composites to Constituent Prop￾erties,” J. Am. Ceram. Soc., 77 [6] 1425–35 (1994). 20E. Vagaggini, J. M. Domergue, and A. G. Evans, “Relationships between Hysteresis Measurements and the Constituent Properties of Ceramic Matrix Compos￾ites: I, Theory,” J. Am. Ceram. Soc., 78 [10] 2709–20 (1995). 21J. M. Domergue, E. Vagaggini, and A. G. Evans, “Relationships between Hysteresis Measurements and the Constituent Properties of Ceramic Matrix Compos￾ites: II, Experimental Studies on Unidirectional Materials,” J. Am. Ceram. Soc., 78 [10] 2721–31 (1995).
32 Journal of the American Ceramic Society—Dong et al. Vol. 86, No. 1