《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_SiC-SiC-4

CERAMICS INTERNATIONAL ELSEⅤIER Ceramics International 28(2002)899-905 www.elsevier.com/locate/ceramint Microstructural evolution and mechanical performances of SiC/SiC composites by polymer impregnation/ microwave pyrolysis (PIMP) process S M. Donga,*, Y. Katoh, A. Kohyama, S.T. Schwab, LL. Snead CREST-ACE, JST and Institute of Advanced Energy, Kyoto University, Gokasho, Uji, Kyoto 611-0011, Japan Thor Technologies, Inc. 7600 Jefferson NE, Suite 9-115. Albuquerque, NM Metals and Ceramic Division, Oak Ridge National Laboratory, PO Box 2008, Oak Ridge, TN 37831-6087, US.A Received 8 January 2002: received in revised form 9 February 2002: accepted 26 March 2002 SiC/SiC composites were prepared by polymer impregnation /microwave pyrolysis(PIMP) process, and their microstructural evolution and the mechanical performances were characterized Using non-coated Tyranno Sa fiber preforms as reinforcement and impregnation with only allylperhydropolycarbosilane(AHPCS) into the preforms, Tyranno SA/SiC composite (TSA/ SiC)with higher density was obtained. While using carbon-coated Tyranno SA fiber preforms, Tyranno SA/C/SiC composite(TSA/C/SiC) h lower density were also fabricated In this composite, Sic particulate was loaded with polymer precursor(AHPCS)in the first cycle impregnation. Microstructural observation revealed that pore and crack formation was affected by processing conditions Bending strength was also dependent on the microstructural evolution of the samples. In TSA/SiC composite, relatively strong interfaces contribute to effective load transfer so that higher bending strength could be reached. In the tsa/C/SiC composite, weak interfaces provide a relatively lower strength. Meanwhile, different microstructural evolution and interfacial properties of the composites lead to the variation of the fracture behaviors. C 2002 Published by elsevier Science Ltd and Techna s.r.I Keywords: B Composites; B. Microstructure-final; C Mechanical properties; D SiC: PIMP 1. Introduction [1-5]. However, because of the lengthy pyrolysis cycles, much time is required in conventional PIP process to Sic/Sic composite is a well-known material promis- produce a dense component, resulting in high-cost for ng for high temperature structural applications because even simple shapes. of its intrinsic thermal stability and excellent mechanical Fast heating, techniques such as microwave or laser properties. Polymer impregnation and pyrolysis(PIP) heating are now being applied for ceramic fabrication or processing is considered to be an effective manufactur- polymer pyrolysis process [6-8]. These novel processing ing technique for preparing high performance SiC/Sic techniques provide a time and energy saving way for composite. Since the shapes of the impregnated parts ceramic preparation. Although those works are still on can be varied intentionally, it will be widely applied to the fundamental stage, promising features have been the complex-shaped components [l] For matrix forma- demonstrated for ceramics development such as short tion, many kinds of polymer precursors have been processing time, uniformity of the products as well as investigated and the details of polymer to ceramic con- tailoring particular design requirements of the materi- version, microstructural development as well as some als. Meanwhile, high frequency microwave radiation physical and mechanical properties have been studied can also be applied to induce the polymer-to-ceramic conversion process, and can induce very high tempera- re in irradiated parts in just minutes [9] The aim of the present study is to characterize the microstructural evolution of SiC/SiC composites 0272-8842/02/$22.00@ 2002 Published by Elsevier Science Ltd and Techna S.r.I PII:S0272-8842(02)00071-8

Microstructural evolution and mechanical performances of SiC/SiC composites by polymer impregnation/microwave pyrolysis (PIMP) process S.M. Donga,*,Y. Katoha ,A. Kohyamaa ,S.T. Schwabb,L.L. Sneadc a CREST-ACE, JST and Institute of Advanced Energy, Kyoto University, Gokasho, Uji, Kyoto 611-0011, Japan bThor Technologies, Inc. 7600 Jefferson NE, Suite 9-115, Albuquerque, NM 87109, USA c Metals and Ceramic Division, Oak Ridge National Laboratory, PO Box 2008, Oak Ridge, TN 37831-6087, USA Received 8 January 2002; received in revised form 9 February 2002; accepted 26 March 2002 Abstract SiC/SiC composites were prepared by polymer impregnation/microwave pyrolysis (PIMP) process,and their microstructural evolution and the mechanical performances were characterized. Using non-coated Tyranno SA fiber preforms as reinforcement and impregnation with only allylperhydropolycarbosilane (AHPCS) into the preforms,Tyranno SA/SiC composite (TSA/SiC) with higher density was obtained. While using carbon-coated Tyranno SA fiber preforms,Tyranno SA/C/SiC composite (TSA/C/SiC) with lower density were also fabricated. In this composite,SiC particulate was loaded with polymer precursor (AHPCS) in the first cycle impregnation. Microstructural observation revealed that pore and crack formation was affected by processing conditions. Bending strength was also dependent on the microstructural evolution of the samples. In TSA/SiC composite,relatively strong interfaces contribute to effective load transfer so that higher bending strength could be reached. In the TSA/C/SiC composite,weak interfaces provide a relatively lower strength. Meanwhile,different microstructural evolution and interfacial properties of the composites lead to the variation of the fracture behaviors. # 2002 Published by Elsevier Science Ltd and Techna S.r.l. Keywords: B. Composites; B. Microstructure-final; C. Mechanical properties; D. SiC; PIMP 1. Introduction SiC/SiC composite is a well-known material promis￾ing for high temperature structural applications because of its intrinsic thermal stability and excellent mechanical properties. Polymer impregnation and pyrolysis (PIP) processing is considered to be an effective manufactur￾ing technique for preparing high performance SiC/SiC composite. Since the shapes of the impregnated parts can be varied intentionally,it will be widely applied to the complex-shaped components [1]. For matrix forma￾tion,many kinds of polymer precursors have been investigated and the details of polymer to ceramic con￾version,microstructural development as well as some physical and mechanical properties have been studied [1–5]. However,because of the lengthy pyrolysis cycles, much time is required in conventional PIP process to produce a dense component,resulting in high-cost for even simple shapes. Fast heating,techniques such as microwave or laser heating are now being applied for ceramic fabrication or polymer pyrolysis process [6–8]. These novel processing techniques provide a time and energy saving way for ceramic preparation. Although those works are still on the fundamental stage,promising features have been demonstrated for ceramics development such as short processing time,uniformity of the products as well as tailoring particular design requirements of the materi￾als. Meanwhile,high frequency microwave radiation can also be applied to induce the polymer-to-ceramic conversion process,and can induce very high tempera￾ture in irradiated parts in just minutes [9]. The aim of the present study is to characterize the microstructural evolution of SiC/SiC composites 0272-8842/02/$22.00 # 2002 Published by Elsevier Science Ltd and Techna S.r.l. PII: S0272-8842(02)00071-8 Ceramics International 28 (2002) 899–905 www.elsevier.com/locate/ceramint * Corresponding author. Tel.: +81-774-38-3465; fax: +81-774-38- 3467. E-mail address: sm-dong@iae.kyoto-u.ac.jp (S.M. Dong)

S.M. Dong et al. Ceramics International 28(2002)899-905 fabricated by polymer impregnation/microwave pyr- all the samples. Fracture surface of each sample after olysis(PIMP) process using 2 D woven Tyranno sa bending test was also observed by sem to determine fiber preforms as the reinforcement. The effects of fiber the fracture behavior. Microstructural evolution of the coating and particulate loading during first cycle AHPCS derived matrix after PIMP process was char impregnation on fiber/matrix interaction and mechan- acterized by transmission electron microscopy (TEM) ical behaviors are also evaluated The thin film was prepared by the focused ion beam (FIB)method. 2. Experimental proced 2.3. Physical and mechanical measurements 2.. Materials All PIMP samples were cut and ground into nearly 4x220 mm rectangular bars to perform three-point Polymer precursor used for matrix formation was bending test at room temperature with an Instron 5581 allylperhydropolycarbosilane(AHPCS), which is regar- test machine. The cross-head speed was 0.5 mm/min ded as a polymer precursor of high ceramic yield [10]. and the span was 18 mm. During bending test, the Since this polymer precursor is very sensitive to the fracture behavior of the bars was in situ recorded by anhydrous and anaerobic environment to minimize the composites was measured by Archimedes's method the ambient atmosphere, it was handled under strictly optical microscope camera. The bulk density of potential for oxygen contamination. The maximum Push-out and push-back tests were performed on thin pyrolysis temperature obtained by microwave irradia- slices of each sample to evaluate the interfacial debond- tion at 37 GHz was approximately 1100oC. The heating ing strength(IDS)and interfacial frictional stress(IFS). time in each cycle was less than 5 min Detailed proces- respectively, using a load controlled micro-indentation sing technique has been described in literature [9] testing system. Those slices were double-face polished Two kinds of SiC/SiC composites were prepared their thickness being less than 100 um to allow the fibers according to PIMP process. The first one is non-coated to be pushed out. During this test, the slices were set on Tyranno SA fiber(Ube Industries Ltd, Japan) rein- a tungsten carbide holder with a 50 um width groove, forced SiC composite (TSA SiC). During impregnation and the load was applied on a single fiber end above the ind pyrolysis, the polymer precursor was used unloaded groove using a triangle diamond indenter. The max- no SiC particulate). Eight cycles of impregnation/pyr- imum load was IN and this load was modified accord- olysis processing were performed. The second is carbon ing to the value of IDS and IFS. The interfacial coated Tyranno SA fiber reinforced Sic composite displacement rate was 0.2 um/s. After push-out test, the (TSA/ C/SiC). In this composite, B-SiC particles(with protruding fibers were firstly observed by SEM, and an average diameter of 0.6 um) were added to the pre- then push-back test was conducted on those protruding cursor in the first cycle of impregnation and then the fibe material was submitted to microwave pyrolysis. This first cycle was followed by five cycles of PIMP with unloaded AHPCS. In the first PIMP cycle, polymer to 3. Results and disscussion filler(SiC) ratio was 50% to 50% by weight. All fiber reinforcements used in this experiment were 2-D woven 3. Microstructural evolution fabrics, and fiber volume fraction of the composites was about 30-35 vol % The fiber volume fraction for com- By contrast to the conventional polymer impregna posite TSA SiC is slightly higher than that in the second tion and pyrolysis(PIP)technique, the PIMP only material. Typical properties of Tyranno SA fiber used in requires very short pyrolysis time. Because the polymer nis experiment are listed in Table I to ceramic conversion is a complex process, such very short pyrolysis duration might affect the evolution of 2. 2. Microstructural observation microstructure. Fig. I shows a typical SEM micrograph Optical microscopy and scanning electron microscopy solidated parts dominate the cross section area. Some (SEM)were conducted on the polished cross-section of isolated large pores could be observed in inter-bundle Table l Properties of Tyranno SA fibers( Grade Im) SiC fiber C/Si atomic ratio Diameter(um) Density (g/cm) Filaments/yarn Tensile strength(GPa) Elastic modulus(GPa) Elongation% Tyranno SA 1.08

fabricated by polymer impregnation/microwave pyr￾olysis (PIMP) process using 2 D woven Tyranno SA fiber preforms as the reinforcement. The effects of fiber coating and particulate loading during first cycle impregnation on fiber/matrix interaction and mechan￾ical behaviors are also evaluated. 2. Experimental procedure 2.1. Materials Polymer precursor used for matrix formation was allylperhydropolycarbosilane (AHPCS),which is regar￾ded as a polymer precursor of high ceramic yield [10]. Since this polymer precursor is very sensitive to the ambient atmosphere,it was handled under strictly anhydrous and anaerobic environment to minimize the potential for oxygen contamination. The maximum pyrolysis temperature obtained by microwave irradia￾tion at 37 GHz was approximately 1100
C. The heating time in each cycle was less than 5 min. Detailed proces￾sing technique has been described in literature [9]. Two kinds of SiC/SiC composites were prepared according to PIMP process. The first one is non-coated Tyranno SA fiber (Ube Industries Ltd.,Japan) rein￾forced SiC composite (TSA/SiC). During impregnation and pyrolysis,the polymer precursor was used unloaded (no SiC particulate). Eight cycles of impregnation/pyr￾olysis processing were performed. The second is carbon coated Tyranno SA fiber reinforced SiC composite (TSA/C/SiC). In this composite, b-SiC particles (with an average diameter of 0.6 mm) were added to the pre￾cursor in the first cycle of impregnation and then the material was submitted to microwave pyrolysis. This first cycle was followed by five cycles of PIMP with unloaded AHPCS. In the first PIMP cycle,polymer to filler (SiC) ratio was 50% to 50% by weight. All fiber reinforcements used in this experiment were 2-D woven fabrics,and fiber volume fraction of the composites was about 30–35 vol.%. The fiber volume fraction for com￾posite TSA/SiC is slightly higher than that in the second material. Typical properties of Tyranno SA fiber used in this experiment are listed in Table 1. 2.2. Microstructural observation Optical microscopy and scanning electron microscopy (SEM) were conducted on the polished cross-section of all the samples. Fracture surface of each sample after bending test was also observed by SEM to determine the fracture behavior. Microstructural evolution of the AHPCS derived matrix after PIMP process was char￾acterized by transmission electron microscopy (TEM). The thin film was prepared by the focused ion beam (FIB) method. 2.3. Physical and mechanical measurements All PIMP samples were cut and ground into nearly 4220 mm rectangular bars to perform three-point bending test at room temperature with an Instron 5581 test machine. The cross-head speed was 0.5 mm/min and the span was 18 mm. During bending test,the fracture behavior of the bars was in situ recorded by optical microscope camera. The bulk density of the composites was measured by Archimedes’s method. Push-out and push-back tests were performed on thin slices of each sample to evaluate the interfacial debond￾ing strength (IDS) and interfacial frictional stress (IFS), respectively,using a load controlled micro-indentation testing system. Those slices were double-face polished, their thickness being less than 100 mm to allow the fibers to be pushed out. During this test,the slices were set on a tungsten carbide holder with a 50 mm width groove, and the load was applied on a single fiber end above the groove using a triangle diamond indenter. The max￾imum load was 1N and this load was modified accord￾ing to the value of IDS and IFS. The interfacial displacement rate was 0.2 mm/s. After push-out test,the protruding fibers were firstly observed by SEM,and then push-back test was conducted on those protruding fibers. 3. Results and disscussion 3.1. Microstructural evolution By contrast to the conventional polymer impregna￾tion and pyrolysis (PIP) technique,the PIMP only requires very short pyrolysis time. Because the polymer to ceramic conversion is a complex process,such very short pyrolysis duration might affect the evolution of microstructure. Fig. 1 shows a typical SEM micrograph of the polished cross-section. Actually,well-con￾solidated parts dominate the cross section area. Some isolated large pores could be observed in inter-bundle Table 1 Properties of Tyranno SA fibers (Grade II) SiC fiber C/Si atomic ratio Diameter (mm) Density (g/cm3 ) Filaments/yarn Tensile strength (GPa) Elastic modulus (GPa) Elongation% Tyranno SA 1.08 10 3.02 800 2.8 420 0.7 900 S.M. Dong et al. / Ceramics International 28 (2002) 899–905

S.M. Dong et al. Ceramics International 28(2002)899-905 age by the pure polymer precursor during the pyrolysis process Microstructure of intra-bundle matrix is shown in Fig. 3. In composite TSA/SiC, as the PIMP process progresses, the matrix became dense. Meanwhile, the initially formed matrix around fibers could also be identified in some areas. The following cycles of PIMP left a clear profile in intra-bundle matrix. Even in this case, the matrix relatively maintained the uniformit around fibers. However, some debonding between fibers and matrix could be observed. In TSA/C/SiC compo site, matrix seems not to be dense even though large amounts of intra-bundle matrix and inter -bundle matrix effectively formed in the first cycle impregnation Fig 1. Typical SEM micrograph on the polished cross-section of the and the particles were not strongly bonded together at composite TSA/SiC the present PIMP temperature. This"porous"matrix might be ascribed to the difficulty for achieving effective polymer impregnation after the matrix was formed in the first cycle impregnation. Generally, during PIMP phenomenon in conventional PIP prepared samples. decrease gradually in size when PIMP cycles proceeded Further observation of the intra-bundle and inter-bun- and then hindered further polymer impregnation. When dle matrix formation was conducted by optical micro- the pores were small enough, the viscous polymer pre- scopy on the polished cross-section, as shown in Fig. 2. cursor could not be effectively impregnated into the Dense intra-bundle matrix can be identified. Meanwhile, consolidated body. At this time, the process should be microcracks still remained but more extensively in the stopped. With particulate loading, the inter-bundle and TSA/C/SiC composite. In this composite, relatively intra-bundle spaces were relatively easier to be filled. large amount of matrix formed in both intra-bundle and After few cycles of PIMP, the remaining porosity, espe inter-bundle areas(Fig. 2b), indicating that the infiltra- cially the open porosity on the outer surface of the tion efficiency was high when particulates were loaded composites was greatly lowered, making the impregna in the first-cycle of impregnation. In composite TSA/ tion more difficult. Since both of the composite densifi SiC (using non-coated Tyranno SA fiber preforms as cation and matrix strengthening are highly dependent reinforcement and without particulate loading during on the polymer impregnation and pyrolysis, insufficient impregnation), intra-bundle fibers and fiber layers were impregnation of polymer precursor implies that the tightly bonded together. Only thin inter-layer matrix matrix might not be strongly bonded In the TSa/C/SiC could be evidenced, as shown by arrow in Fig 2a. These composite, polymer impregnation was stopped after six results might be ascribed to low mass inclusion( without cycles of PIMP, while for TSA/SiC, eight cycles could the addition of Sic particles) and large volume shrink- Fig. 2. Optical photographs of the polished cross-section of the composites showing the intra-bundle and inter-bundle matrix formation and the cracks propagation: (a) TSA/SiC, (b) TSA/C/SiC

areas and inter-layers. This is also the typically observed phenomenon in conventional PIP prepared samples. Further observation of the intra-bundle and inter-bun￾dle matrix formation was conducted by optical micro￾scopy on the polished cross-section,as shown in Fig. 2. Dense intra-bundle matrix can be identified. Meanwhile, microcracks still remained but more extensively in the TSA/C/SiC composite. In this composite,relatively large amount of matrix formed in both intra-bundle and inter-bundle areas (Fig. 2b),indicating that the infiltra￾tion efficiency was high when particulates were loaded in the first-cycle of impregnation. In composite TSA/ SiC (using non-coated Tyranno SA fiber preforms as reinforcement and without particulate loading during impregnation),intra-bundle fibers and fiber layers were tightly bonded together. Only thin inter-layer matrix could be evidenced,as shown by arrow in Fig. 2a. These results might be ascribed to low mass inclusion (without the addition of SiC particles) and large volume shrink￾age by the pure polymer precursor during the pyrolysis process. Microstructure of intra-bundle matrix is shown in Fig. 3. In composite TSA/SiC,as the PIMP process progresses,the matrix became dense. Meanwhile,the initially formed matrix around fibers could also be identified in some areas. The following cycles of PIMP left a clear profile in intra-bundle matrix. Even in this case,the matrix relatively maintained the uniformity around fibers. However,some debonding between fibers and matrix could be observed. In TSA/C/SiC compo￾site,matrix seems not to be dense even though large amounts of intra-bundle matrix and inter-bundle matrix were effectively formed in the first cycle impregnation, and the particles were not strongly bonded together at the present PIMP temperature. This ‘‘porous’’ matrix might be ascribed to the difficulty for achieving effective polymer impregnation after the matrix was formed in the first cycle impregnation. Generally,during PIMP process,the micropores left in the matrix would decrease gradually in size when PIMP cycles proceeded and then hindered further polymer impregnation. When the pores were small enough,the viscous polymer pre￾cursor could not be effectively impregnated into the consolidated body. At this time,the process should be stopped. With particulate loading,the inter-bundle and intra-bundle spaces were relatively easier to be filled. After few cycles of PIMP,the remaining porosity,espe￾cially the open porosity on the outer surface of the composites was greatly lowered,making the impregna￾tion more difficult. Since both of the composite densifi- cation and matrix strengthening are highly dependent on the polymer impregnation and pyrolysis,insufficient impregnation of polymer precursor implies that the matrix might not be strongly bonded. In the TSA/C/SiC composite,polymer impregnation was stopped after six cycles of PIMP,while for TSA/SiC,eight cycles could be performed. Fig. 1. Typical SEM micrograph on the polished cross-section of the composite TSA/SiC. Fig. 2. Optical photographs of the polished cross-section of the composites showing the intra-bundle and inter-bundle matrix formation and the cracks propagation: (a) TSA/SiC,(b) TSA/C/SiC. S.M. Dong et al. / Ceramics International 28 (2002) 899–905 901

S.M. Dong et al. Ceramics International 28(2002)899-905 Fig 3. Comparison of the microstructural evolution of the PIMP composites:(a),(b)TSA/SiC;(c).(d) TSA/C/SiC. The microstructure of the composites at a higher 3. 2. Physical and mechanical properties magnification is shown in Fig. 3b and d. Some micro- cracks can be observed around the fibers. Even though Table 2 lists some physical and mechanical properties large amounts of inter-bundle and intra-bundle matrix of the composites. Bulk density of TSA SiC composite were effectively formed in TSA/C/SiC composite, the is relatively high(2.51 g/cm)compared to that(2.30 g/ intra-bundle matrix seems to be loosely consolidated. cm) for TSA/C/SiC composite In composite TSA/SiC. This feature could be distinguished from the inter-bun- as discussed in the previous section, intra-bundle matrix dle matrix areas was gradually grown around the fibers when increasing TEM observation of the matrix from AHPCS pre- the number of impregnation cycles. After eight cycles cursor is shown in Fig. 4. It can be evidenced from the PIMP process, polymer derived matrix was formed in SAD-pattern that the matrix is completely amorphous, the intra-bundle and inter-bundle areas. Although some probably due to the very short pyrolysis time and low big pores inevitably existed, the density of TSA/SiC pyrolysis temperature(around 1100C) could still reach 2.51 g/cm. In the TSA/C/SiC compo site, the formation of relatively loose microstructure with the inclusion of Sic particulates and the les efficient impregnation cycles might explain the lower Fiber density. Bending test results indicate that higher strength Interface could be obtained for the composite with higher density (TSA/SiC). The average strength of this composite is over 400 MPa. While for the composites with lower density, strength is at a lower level Matrix To better understand the interaction between fibers and matrix, push-out and push-back tests were con- ducted on each composite. Typical curves are shown in 500n Fig. 5a. Protruding fiber is also demonstrated as an example in Fig 5b. In TSA/C/SiC, debonding mainly Fig. 4. TEM image and SAD-pattern (inset) of the composite TSA/ occurred in the fiber/carbon interface during push-out Sic showing the formation of non-crystallized matrix. test. In Fig. 5, Po represents the push-out load, which

The microstructure of the composites at a higher magnification is shown in Fig. 3b and d. Some micro￾cracks can be observed around the fibers. Even though large amounts of inter-bundle and intra-bundle matrix were effectively formed in TSA/C/SiC composite,the intra-bundle matrix seems to be loosely consolidated. This feature could be distinguished from the inter-bun￾dle matrix areas. TEM observation of the matrix from AHPCS pre￾cursor is shown in Fig. 4. It can be evidenced from the SAD-pattern that the matrix is completely amorphous, probably due to the very short pyrolysis time and low pyrolysis temperature (around 1100
C). 3.2. Physical and mechanical properties Table 2 lists some physical and mechanical properties of the composites. Bulk density of TSA/SiC composite is relatively high (2.51 g/cm3 ) compared to that (2.30 g/ cm3 ) for TSA/C/SiC composite. In composite TSA/SiC, as discussed in the previous section,intra-bundle matrix was gradually grown around the fibers when increasing the number of impregnation cycles. After eight cycles PIMP process,polymer derived matrix was formed in the intra-bundle and inter-bundle areas. Although some big pores inevitably existed,the density of TSA/SiC could still reach 2.51 g/cm3 . In the TSA/C/SiC compo￾site,the formation of relatively loose microstructure with the inclusion of SiC particulates and the less efficient impregnation cycles might explain the lower density. Bending test results indicate that higher strength could be obtained for the composite with higher density (TSA/SiC). The average strength of this composite is over 400 MPa. While for the composites with lower density,strength is at a lower level. To better understand the interaction between fibers and matrix,push-out and push-back tests were con￾ducted on each composite. Typical curves are shown in Fig. 5a. Protruding fiber is also demonstrated as an example in Fig. 5b. In TSA/C/SiC,debonding mainly occurred in the fiber/carbon interface during push-out test. In Fig. 5, Po represents the push-out load,which Fig. 3. Comparison of the microstructural evolution of the PIMP composites: (a),(b) TSA/SiC; (c),(d) TSA/C/SiC. Fig. 4. TEM image and SAD-pattern (inset) of the composite TSA/ SiC showing the formation of non-crystallized matrix. 902 S.M. Dong et al. / Ceramics International 28 (2002) 899–905

S.M. Dong et al./ Ceramics International 28(2002 )899-905 TSA/SiC. Although microcracks around some fibers Push-out·Push have been formed, the average IDS still reaches 161.1 MPa. In composites TSA/C/SiC, relatively weak inter- faces were formed. The data for interfacial frictional stress reveal the same trend as that of the IDs. TSA/SiC composite also shows a higher IFS value. However, it should be mentioned that debonding between fibers and matrix could still be realized under certain load during push-out test in TSA SiC although the matrix and fibers were directly bonded during the PIMP 00 The differences of interfacial debonding strength and interfacial frictional stress might be ascribed to the interaction between fibers and matrix [12-15. Gen- erally, the radial stress on the fibers can be described by where o is radial thermal residual stress given by ar=k(am-a)△T (3) or is fiber surface roughness induced radial misfit stress given by Fig. 5. Typical single fiber push-out and push back load/ displacem curves(a), and SEM micrographs of the protruding fibers after singl (4) fiber push-out test(b) for composite TSA/ SiC. wEmer means the load for fiber debonding and sliding as well Em(l-ur)+ er(I Um) as the contributions from fiber elastic/plastic deforma- tion and Poisson expansion effects [ll]. It is used in this oP is the radial stress induced by Poisson's effect. ar study for estimating the interfacial debonding strength, And af are the thermal expansion coefficients of the IDS via the formula fiber and matrix in the radial direction. AT is the tem- IDS= Po/IDt () perature difference between composite fabrication tem- perature (1373 K)and room temperature (298 K),at which the push-out and push-back test were conducted where D and t are the fiber diameter and the specimen Er, Em, Uf and umare Youngs moduli and Poisson,s thickness of the slice for push-out and push-back test. ratios of the fiber and matrix. 8 Is the fiber surface The same formula was also used for determining the roughness. In this experiment, fiber surface roughness interfacial frictional stress(IFS), where Po was sub- profiles for Tyranno SA fiber was examined by an opti- stituted by Pb, which is mainly used to initiate the slid al interferometric microscope, MicromapTM, with of resolution of less than 1 nm in z-direction and 8 was From the calculated results listed in Table 2. it can be determined to be 3. 74. o Is a factor for the effect of fiber bserved that a relatively strong interface is present in volume fraction in the composites. Here, oP and o are simply assumed to be zero and l, respectively. Since the Table 2 Tyranno SA is well crystallized and near stoichiometric Physical and mechanical properties of the composites [17], their thermal expansion coefficients can be assumed to be similar as those of the sintered sic monolithic Composites ISA/SIC TSA/C/SIC ceramics(about 4.0x10-6/K)[18]. For the polymer Bulk density (g/cm) 2.51士0.0 2.30±0.04 derived matrix, thermal expansion coefficient can be Flexural strength(MPa) 402.4*35.9 258.7-+26.0 simply assumed to be 3.0-3.5x10-6/K(similar to that Interfacial debonding strength(MPa) 161.1+31.6 61.5+17.8 of a Si-C[O] amorphous phase)[18, 19]. Based on those Interfacial frictional stress(MPa) 56.9±18.2 Modulus of elasticity(GPa) 1020±5.3 49.7±4.8 assumptions, the radial stress on the fibers in the com- posites can be roughly estimated using Eqs. (2H5)

means the load for fiber debonding and sliding as well as the contributions from fiber elastic/plastic deforma￾tion and Poisson expansion effects [11]. It is used in this study for estimating the interfacial debonding strength, IDS via the formula: IDS ¼ Po=Dt ð1Þ where D and t are the fiber diameter and the specimen thickness of the slice for push-out and push-back test. The same formula was also used for determining the interfacial frictional stress (IFS),where Po was sub￾stituted by Pb,which is mainly used to initiate the slid￾ing of the debonded fiber. From the calculated results listed in Table 2,it can be observed that a relatively strong interface is present in TSA/SiC. Although microcracks around some fibers have been formed,the average IDS still reaches 161.1 MPa. In composites TSA/C/SiC,relatively weak inter￾faces were formed. The data for interfacial frictional stress reveal the same trend as that of the IDS. TSA/SiC composite also shows a higher IFS value. However,it should be mentioned that,debonding between fibers and matrix could still be realized under certain load during push-out test in TSA/SiC although the matrix and fibers were directly bonded during the PIMP process. The differences of interfacial debonding strength and interfacial frictional stress might be ascribed to the interaction between fibers and matrix [12–15]. Gen￾erally,the radial stress on the fibers can be described by [16], N r ¼ T r þ R r þ P r ð2Þ where T r is radial thermal residual stress given by T r ¼  r m  r f
T ð3Þ R r is fiber surface roughness induced radial misfit stress given by R r ¼   Rf  ð4Þ  ¼ !EmEf Emð Þþ 1  f Ef ð Þ 1 þ m ð5Þ P r is the radial stress induced by Poisson’s effect. r m And r f are the thermal expansion coefficients of the fiber and matrix in the radial direction. T is the tem￾perature difference between composite fabrication tem￾perature (1373 K) and room temperature (298 K),at which the push-out and push-back test were conducted. Ef, Em, r f and r mare Young’s moduli and Poisson’s ratios of the fiber and matrix.  Is the fiber surface roughness. In this experiment,fiber surface roughness profiles for Tyranno SA fiber was examined by an opti￾cal interferometric microscope,MicromapTM,with a resolution of less than 1 nm in Z-direction,and  was determined to be 3.74. ! Is a factor for the effect of fiber volume fraction in the composites. Here, P r and ! are simply assumed to be zero and 1,respectively. Since the Tyranno SA is well crystallized and near stoichiometric [17],their thermal expansion coefficients can be assumed to be similar as those of the sintered SiC monolithic ceramics (about 4.0106 /K) [18]. For the polymer derived matrix,thermal expansion coefficient can be simply assumed to be 3.0–3.5106 /K (similar to that of a Si–C[–O] amorphous phase) [18,19]. Based on those assumptions,the radial stress on the fibers in the com￾posites can be roughly estimated using Eqs. (2)–(5). Table 2 Physical and mechanical properties of the composites Composites TSA/SiC TSA/C/SiC Bulk density (g/cm3 ) 2.510.02 2.300.04 Flexural strength (MPa) 402.435.9 258.726.0 Interfacial debonding strength (MPa) 161.131.6 61.517.8 Interfacial frictional stress (MPa) 56.918.2 32.810.1 Modulus of elasticity (GPa) 102.05.3 49.74.8 Fig. 5. Typical single fiber push-out and push back load/displacement curves (a),and SEM micrographs of the protruding fibers after single fiber push-out test (b) for composite TSA/SiC. S.M. Dong et al. / Ceramics International 28 (2002) 899–905 903

S.M. Dong et al. Ceramics International 28(2002)899-905 curve of TSA/C/SiC demonstrates a different behavior with respect to the TSA/SiC composite. After reaching TSA/S the maximum value, the load decreases gradually indi- cating a pseudo-ductile fracture behavior. This char acteristic might be ascribed to the relatively loose matrix and weak interface also, providing the lower elastic TSA/C/SIC modulus. Even though the weak interface is benefitial for the crack bridging and fiber pull-out, it is simulta neously detrimental for strength because of the low load transfer ability from matrix to fibers through the weak interface Fig. 6. Stress/displacement curves of the composites. 3.3. Fracture behaviors of the composites Because am af in the composites, a radial tension Fig. 7 shows the fracture surfaces of the composites stress was induced to the fiber/matrix interface by ther- For TSA/SiC, short fiber pull-out is mal mismatch. This might explain the occurrence of Although polymer derived matrix directly bonded with some fiber/matrix debonding as indicated in Fig 3 For fibers, this bonding was weakened by the tensile effect TSA/C/SiC, the carbon coating will simultaneously induced by thermal mismatch so that the debonding affect the interfacial stresses [17]. The radial tension between fibers and matrix became possible. In situ stress can weaken the fiber/matrix interfacial bonding observation during bending test indicated that crack and significantly reduce the applied load during the fiber propagated nearly parallel to the direction of the load push-out and push-back [13]. This effect might be bene- applied. No delamination along fiber layers could be ficial for the composite TSA SiC to prevent a too strong found and the tensile fracture is predominant. TSA/ CI fiber/matrix bonding in the case of no coating applied. Sic composite demonstrates long fiber pull-out, as The roughness induced radial stress will highly con- shown in Fig. 7b. The in situ observation indicated that ribute to the IFS, as analyzed in literature [13, 14, 16]. In the fracture propagated in two directions: one is nearly the present study, Tyranno SA fiber is confirmed to parallel to that of the applied load and another is nearly have a rather rough surface so that the contribution to erpendicular to that of the applied load. It means that the IFS is high. This characteristic will finally affect the both tensile fracture and shear fracture exist in this mechanical properties of the composites. composite Typical stress/displacement curves obtained from The dense matrix in the TSA SiC composite would bending test are shown in Fig. 6. TSA/SiC composite prevent delamination along the fiber layers. The strong displays not only higher strength but also a high mod- bonding was beneficial for load transfer from matrix to ulus of elasticity, as summarized in Table 2. These fibers and simultaneously limited the fibers pull-out. In mechanical properties may be attributed to the densely TSA/C/SiC composite, since the density is relatively low formed matrix and the strong interface between fibers and the matrix was loosely formed during PIMP pro- and matrix. The strong interface is beneficial for the cess, cracks easily propagate along the weak region load transfer from matrix to fibers so that higher Meanwhile, the weak load transfer between matrix and strength could be obtained [20, 21]. Load/ displacement fibers in this composite allows the long fibers pull-out SOum Fig. 7. SEM micrographs of the fracture surface after bending test showing the different behaviors of the PIMP composites: (a) TSA/SiC, (b) TSA

Because r m < r f in the composites,a radial tension stress was induced to the fiber/matrix interface by ther￾mal mismatch. This might explain the occurrence of some fiber/matrix debonding as indicated in Fig. 3. For TSA/C/SiC,the carbon coating will simultaneously affect the interfacial stresses [17]. The radial tension stress can weaken the fiber/matrix interfacial bonding and significantly reduce the applied load during the fiber push-out and push-back [13]. This effect might be bene- ficial for the composite TSA/SiC to prevent a too strong fiber/matrix bonding in the case of no coating applied. The roughness induced radial stress will highly con￾tribute to the IFS, as analyzed in literature [13,14,16]. In the present study,Tyranno SA fiber is confirmed to have a rather rough surface so that the contribution to the IFS is high. This characteristic will finally affect the mechanical properties of the composites. Typical stress/displacement curves obtained from bending test are shown in Fig. 6. TSA/SiC composite displays not only higher strength but also a high mod￾ulus of elasticity,as summarized in Table 2. These mechanical properties may be attributed to the densely formed matrix and the strong interface between fibers and matrix. The strong interface is beneficial for the load transfer from matrix to fibers so that higher strength could be obtained [20,21]. Load/displacement curve of TSA/C/SiC demonstrates a different behavior with respect to the TSA/SiC composite. After reaching the maximum value,the load decreases gradually indi￾cating a pseudo-ductile fracture behavior. This char￾acteristic might be ascribed to the relatively loose matrix and weak interface also,providing the lower elastic modulus. Even though the weak interface is benefitial for the crack bridging and fiber pull-out,it is simulta￾neously detrimental for strength because of the low load transfer ability from matrix to fibers through the weak interface. 3.3. Fracture behaviors of the composites Fig. 7 shows the fracture surfaces of the composites. For TSA/SiC,short fiber pull-out is apparent (Fig. 7a). Although polymer derived matrix directly bonded with fibers,this bonding was weakened by the tensile effect induced by thermal mismatch so that the debonding between fibers and matrix became possible. In situ observation during bending test indicated that crack propagated nearly parallel to the direction of the load applied. No delamination along fiber layers could be found,and the tensile fracture is predominant. TSA/C/ SiC composite demonstrates long fiber pull-out,as shown in Fig. 7b. The in situ observation indicated that the fracture propagated in two directions: one is nearly parallel to that of the applied load and another is nearly perpendicular to that of the applied load. It means that both tensile fracture and shear fracture exist in this composite. The dense matrix in the TSA/SiC composite would prevent delamination along the fiber layers. The strong bonding was beneficial for load transfer from matrix to fibers and simultaneously limited the fibers pull-out. In TSA/C/SiC composite,since the density is relatively low and the matrix was loosely formed during PIMP pro￾cess,cracks easily propagate along the weak region. Meanwhile,the weak load transfer between matrix and fibers in this composite allows the long fibers pull-out. Fig. 6. Stress/displacement curves of the composites. Fig. 7. SEM micrographs of the fracture surface after bending test showing the different behaviors of the PIMP composites: (a) TSA/SiC,(b) TSA/ C/SiC. 904 S.M. Dong et al. / Ceramics International 28 (2002) 899–905

S.M. Dong et al. Ceramics International 28(2002)899-905 4. Conclusions [6 HH. Strecker, K P. Norton, J D. Katz, J.O. Freim, Microwave densification of electrophoretically infiltrated silicon carbide The composite without carbon coating and particu composite, J Mater. Sci. 32(1997)6429-6433 late loading demonstrates higher density and flexural [7 B.G. Ravi, V. Praveen, M. Panneer Sevam, K.J. Rao, Micro- strength. In this composite, a relative strong interface ave-assisted preparation and sintering of mullite and mullite. zirconia composites from metal organics, Mater. Res. Bull. 33 between matrix and fibers could be formed. The strong (1998)1527-1536 interface can provide the high ability for load transfer [8K. Jakubenas, H L. Marcus, Silicon carbide from laser pyrolysis from matrix to fibers and prevent the delamination of of polycarbosilane, J. Am. Ceram Soc. 78(1995)2263-2266 he fiber layers 9 S.T. Schwab, P F. Fleig T Chen, J D. Katz, T.w. Hardek, K w Buesking, Enhanced PIP processing of Sic/SiC for fusion appli The composites with carbon coating and particulate cations, in: Proceeding 3rd IEA International Workshop on Si loading present a relatively lower density and flexural Sic Ceramic Composites for Fusion Applications, 1998 strength. This composite has a relatively weak interface [] L.v. Interrante, C.W.Whitmarsh, w. Sherwood, H.Wu nd shows long fiber pull-out. During bending test shear fracture occurred leading to the delamination Res. Soc. Symp Proc. 346(1994)593-603 along the interlayers. With the adjustment of the filler [11 T. Hinoki, w. Yang, T. Nozawa, T. Shibayama, Y. Katoh addition in the impregnated polymer precursor and fiber A. Kohyama, Improvement of mechanical properties of SiC/sic coating, interfacial properties could be changed. This composites by various surface treatments of fibers, J. Nucl. implies that the fracture behaviors can be further mod Mater.289(2001)2329 ified to fit the requirement for preparing strong compo- [2]RJ. Kerans, R.S. Hay, N.J. Pagano, T.A. Parthasarathy, The role of the fiber-matrix interface in ceramic sites and or tough composites by PIMP process. Bul.68(1989)429442. [13 P D Jero, R.J. Kerans, T.A. Parthasarathy, Effect of interfacial roughness on the frictional stress measured using pushout tests, J Acknowledgements Am. Ceram.Soc.74(1991)2793-2801. [14 T.A. Parthasarathy, D B. Marshall, RJ. Kerans, This work was conducted at the Institute of the effect of interfacial roughness of fiber debonding and sti ing rittle matrix composites, Acta Metall. Mater. 42(1994)3773- Advanced Energy, Kyoto University, as a part of Core Research for Evolutional Science and Technology [15] F. Rebillat, J. Lamon, R. Naslain, E L. Curzio, M K Fe (CREST) program administered by Japan Science and T M. Besmann, Properties of multilayered interphases in SiC/SiC Technology Corporation JST) interfaces, J Am Ceram Soc. 81(1998)965-978. [16Y. Tanaka, Y Kagawa, Y.F. Liu, C Masuda, mechanism during high temperature fatigue References reinforced Ti alloy matrix composite, Mater. (2001)110-117 [K. Sato, A. Tezuka, O. Funayama, T. Isoda, Y. Terada, S. Kato, [17 S.M. Dong, G. Chollon, C. Labrugere, M. Lahaye, A. Guette, M. Iwata, Fabrication and pressure testing of a gas-turbine R. Naslain. D L Jiang, Characterization of some advanced Si based ceramic fibers, J. Mater. Sci. 36(2001) method, Comp. Sci. Tech 59(1999)853-8 [18 J.L. Bobet, J. Lamon, Thermal residual stress [2 G.D. Soraru, F. Babonneau, J.D. Mackenzie, Structural evolu- mposites-1 Axisymmetrical model and finite tions from polycarbosilane to SiC ceramics, J. Mater. Sci. 25 Acta Metall. Mater. 43(1995)2241-2253 (1990)3886-3893 [19 C.A. Hasegawa, A. Kohyama, R.J. Jones, L L. Snead, B] MJ. Wild, P. Buhler, On the phase composition of poly- B. Riccardi, P. Fenici, Critical issues and current status of Sic/ nethylsiloxane derived ceramics, J. Mater. Sci. 33(1998)5441 Sic composites for fusion, J Nucl. Mater. 283-287(2000)128- 544 4 G. Ziegler, I. Richter, D. Suttor, Fiber-reinforced composites [20 F. Rebillat, J. Lamon, A Guette, The concept of a strong inter- with polymer-derived matrix: processing, matrix formation and face applied to SiC/SiC composites with a BN interphase. Acta properties, Composites A30(1999)411-417. Mater.48(2000)4609-4618 5 M. Kotani, A. Kohyama, K. Okamuram, T. Inoue, Fabrication [21] F. Rebillat, J. Lamon, R. Naslain, E L. Curzio, M K. Ferber, f high performance SiC/SiC compos site b polymer Impregn T M. Besmann, Interfacial bond strength in SiC/SiC composite tion and pyrolysis method, Ceram. Eng. Sci. Proc. 20(1999)309- by single-fiber push-out tests, J Am Ceram. 316. Soc.8l(1998)2315-2326

4. Conclusions The composite without carbon coating and particu￾late loading demonstrates higher density and flexural strength. In this composite,a relative strong interface between matrix and fibers could be formed. The strong interface can provide the high ability for load transfer from matrix to fibers and prevent the delamination of the fiber layers. The composites with carbon coating and particulate loading present a relatively lower density and flexural strength. This composite has a relatively weak interface and shows long fiber pull-out. During bending test, shear fracture occurred leading to the delamination along the interlayers. With the adjustment of the filler addition in the impregnated polymer precursor and fiber coating,interfacial properties could be changed. This implies that the fracture behaviors can be further mod￾ified to fit the requirement for preparing strong compo￾sites and/or tough composites by PIMP process. Acknowledgements This work was conducted at the Institute of the Advanced Energy,Kyoto University,as a part of Core Research for Evolutional Science and Technology (CREST) program administered by Japan Science and Technology Corporation (JST). References [1] K. Sato,A. Tezuka,O. Funayama,T. Isoda,Y. Terada,S. Kato, M. Iwata,Fabrication and pressure testing of a gas-turbine component manufactured by a preceramic-polymer-impregnation method,Comp. Sci. Tech. 59 (1999) 853–859. [2] G.D. Soraru,F. Babonneau,J.D. Mackenzie,Structural evolu￾tions from polycarbosilane to SiC ceramics,J. Mater. Sci. 25 (1990) 3886–3893. [3] M.J. Wild,P. Buhler,On the phase composition of poly￾methylsiloxane derived ceramics,J. Mater. Sci. 33 (1998) 5441– 5444. [4] G. Ziegler,I. Richter,D. Suttor,Fiber-reinforced composites with polymer-derived matrix: processing,matrix formation and properties,Composites A30 (1999) 411–417. [5] M. Kotani,A. Kohyama,K. Okamuram,T. Inoue,Fabrication of high performance SiC/SiC composite by polymer impregna￾tion and pyrolysis method,Ceram. Eng. Sci. Proc. 20 (1999) 309– 316. [6] H.H. Streckert,K.P. Norton,J.D. Katz,J.O. Freim,Microwave densification of electrophoretically infiltrated silicon carbide composite,J. Mater. Sci. 32 (1997) 6429–6433. [7] B.G. Ravi,V. Praveen,M. Panneer Sevam,K.J. Rao,Micro￾wave-assisted preparation and sintering of mullite and mullite￾zirconia composites from metal organics,Mater. Res. Bull. 33 (1998) 1527–1536. [8] K. Jakubenas,H.L. Marcus,Silicon carbide from laser pyrolysis of polycarbosilane,J. Am. Ceram. Soc. 78 (1995) 2263–2266. [9] S.T. Schwab,P.F. Fleig,T. Chen,J.D. Katz,T.W. Hardek,K.W. Buesking,Enhanced PIP processing of SiC/SiC for fusion appli￾cations,in: Proceeding 3rd IEA International Workshop on SiC/ SiC Ceramic Composites for Fusion Applications,1998. [10] L.V. Interrante,C.W. Whitmarsh,W. Sherwood,H.J. Wu, R. Lewis,G. Maciel,High yield polycarbosilane precursors to stoichiometric SiC: synthesis,pyrolysis and application,Mater. Res. Soc. Symp. Proc. 346 (1994) 593–603. [11] T. Hinoki,W. Yang,T. Nozawa,T. Shibayama,Y. Katoh, A. Kohyama,Improvement of mechanical properties of SiC/SiC composites by various surface treatments of fibers,J. Nucl. Mater. 289 (2001) 23–29. [12] R.J. Kerans,R.S. Hay,N.J. Pagano,T.A. Parthasarathy,The role of the fiber–matrix interface in ceramic composites,Ceram. Bull. 68 (1989) 429–442. [13] P.D. Jero,R.J. Kerans,T.A. Parthasarathy,Effect of interfacial roughness on the frictional stress measured using pushout tests,J. Am. Ceram. Soc. 74 (1991) 2793–2801. [14] T.A. Parthasarathy,D.B. Marshall,R.J. Kerans,Analysis of the effect of interfacial roughness of fiber debonding and sliding in brittle matrix composites,Acta Metall. Mater. 42 (1994) 3773– 3784. [15] F. Rebillat,J. Lamon,R. Naslain,E.L. Curzio,M.K. Ferber, T.M. Besmann,Properties of multilayered interphases in SiC/SiC chemical-vapor-infiltrated composites with ‘‘weak’’ and ‘‘strong’’ interfaces,J. Am. Ceram. Soc. 81 (1998) 965–978. [16] Y. Tanaka,Y. Kagawa,Y.F. Liu,C. Masuda,Interface damage mechanism during high temperature fatigue test in SiC fiber￾reinforced Ti alloy matrix composite,Mater. Sci. Eng. A314 (2001) 110–117. [17] S.M. Dong,G. Chollon,C. Labrugere,M. Lahaye,A. Guette, R. Naslain,D.L. Jiang,Characterization of some advanced SiC￾based ceramic fibers,J. Mater. Sci. 36 (2001) 2371–2381. [18] J.L. Bobet,J. Lamon,Thermal residual stress in ceramic matrix composites—I. Axisymmetrical model and finite element analysis, Acta Metall. Mater. 43 (1995) 2241–2253. [19] C.A. Hasegawa,A. Kohyama,R.J. Jones,L.L. Snead, B. Riccardi,P. Fenici,Critical issues and current status of SiC/ SiC composites for fusion,J. Nucl. Mater. 283–287 (2000) 128– 137. [20] F. Rebillat,J. Lamon,A. Guette,The concept of a strong inter￾face applied to SiC/SiC composites with a BN interphase,Acta Mater. 48 (2000) 4609–4618. [21] F. Rebillat,J. Lamon,R. Naslain,E.L. Curzio,M.K. Ferber, T.M. Besmann,Interfacial bond strength in SiC/SiC composite materials,as studied by single-fiber push-out tests,J. Am. Ceram. Soc. 81 (1998) 2315–2326. S.M. Dong et al. / Ceramics International 28 (2002) 899–905 905