《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_C-SiC-7

MATERIALS HIENGE& ENGIEERING ELSEVIER Materials Science and Engineering A 465 (2007)290-294 www.elsevier.com/locate/msea Short communication Fabrication and characterization of C/Sic composites with large thickness, high density and near-stoichiometric matrix by heaterless chemical vapor infiltration Sufang Tang, Jingyi Deng, Shijun Wang, Wenchuan Liu Instinte of Metal Research, Chinese Academy of Sciences, Shenyang 110016, PR China Received 14 December 2006: received in revised form I February 2007; accepted 6 February 2007 The C/SiC composites with a 7-8 mm thickness, a 2.3-2.4 g/em'density, and a Si/C ratio 1.08-1.09 matrix have been fabricated by heaterless chemical vapor infiltration(CVi) in a 25-h deposition time. The C/siC composites exhibited a damage-tolerant fracture behavior, and its average flexural strength, longitudinal/transverse compressive strength and fracture toughness were 163 MPa, 304/276 MPa and 6.5 MPam, respectively The composites presented a thermal expansion of (1.09-1.93)10C- and a thermal conductivity (tC) of 8.34-6.56 W/m"C in the range 200-1000°C 2007 Elsevier B. v. All rights reserved. Keywords: C/SiC composites: CVI; Mechanical properties; Thermal conductivity 1. Introduction the densification process, and also because large numbers of arying preforms are easily accommodated in a reactor [3-5] As material requirements become more and more sophis- However, the process has several disadvantages for fabricating ticated, there seems to be a substantially greater emphasis on the composites: low deposition rate, time-consuming treatment, development of advanced thermo-structural composites such as severe corrosion of vacuum pump, and consequent high cost C/C, C/SiC and SiC/SiC, because of their low density, high- The liquid silicon infiltration(LsD) is an attractive method for temperature strength, thermal shock resistance, low coefficient preparation SiC composites due to an effective cost of thermal expansion(CTE), good thermal conductivity (TC), [6-8. The method is infiltration of molten Si into low den- abrasion resistance and high fracture toughness. The excellent sity C/C and transformation by a reactive sintering process into mechanical and thermo-physical properties provide many poten- the dense materials. Nevertheless, the residual Si metal and the tial applications for these materials, including re-entry shields, etched fibers will result in poor high-temperature strength of the rocket nozzles, disc brakes, etc. [1, 2]. However, the major obsta- composites. Another interesting method is the liquid polymer cle to a much broader application in the civil sector, such as disc impregnation(LPD)method based on organometallic precur- brakes, heat exchangers, gas turbines and chemical reactors, is sors such as polycarbosilane [9-11]. However, high numbers till their quite high price of impregnation/pyrolysis steps employed for manufacturing a Several main processes, including chemical vapor infiltra- relatively dense material are required because of the high poros- tion(CVi), liquid and solid infiltration and polymer conversion, ity of the initial fiber preform and the low ceramic yield of the have been employed for fabrication of this new class of the precursor materials 3-11]. The isothermal CVi method is considered as Therefore, it is required to develop a low-cost process for one of the most promising techniques because its major advan- fabricating the high-performance thermo-structural composites tage is the lower thermal and mechanical stress achieved during Such an attempt has been made and a novel low-cost manufac- turing method has been developed, primarily for of the C/C, C/SiC and SiC/SiC composites [12, 13]. In this Corresponding author. Tel. +86 2483978056 paper, large thickness, high density and near-stoichiometric E-mail address: jydeng @imr ac cn ( Deng) matrix C/SiC composites were fabricated by a simple heaterless 0921-5093/S-see front matter O 2007 Elsevier B v. All rights reserved doi:10.1016/msea.2007.02037

Materials Science and Engineering A 465 (2007) 290–294 Short communication Fabrication and characterization of C/SiC composites with large thickness, high density and near-stoichiometric matrix by heaterless chemical vapor infiltration Sufang Tang, Jingyi Deng ∗, Shijun Wang, Wenchuan Liu Institute of Metal Research, Chinese Academy of Sciences, Shenyang 110016, PR China Received 14 December 2006; received in revised form 1 February 2007; accepted 6 February 2007 Abstract The C/SiC composites with a 7–8 mm thickness, a 2.3–2.4 g/cm3 density, and a Si/C ratio 1.08–1.09 matrix have been fabricated by heaterless chemical vapor infiltration (CVI) in a 25-h deposition time. The C/SiC composites exhibited a damage-tolerant fracture behavior, and its average flexural strength, longitudinal/transverse compressive strength and fracture toughness were 163 MPa, 304/276 MPa and 6.5 MPa m1/2, respectively. The composites presented a thermal expansion of (1.09–1.93) × 10−6 ◦C−1 and a thermal conductivity (TC) of 8.34–6.56 W/m ◦C in the range 200–1000 ◦C. © 2007 Elsevier B.V. All rights reserved. Keywords: C/SiC composites; CVI; Mechanical properties; Thermal conductivity 1. Introduction As material requirements become more and more sophis￾ticated, there seems to be a substantially greater emphasis on development of advanced thermo-structural composites such as C/C, C/SiC and SiC/SiC, because of their low density, high￾temperature strength, thermal shock resistance, low coefficient of thermal expansion (CTE), good thermal conductivity (TC), abrasion resistance and high fracture toughness. The excellent mechanical and thermo-physical properties provide many poten￾tial applications for these materials, including re-entry shields, rocket nozzles, disc brakes, etc. [1,2]. However, the major obsta￾cle to a much broader application in the civil sector, such as disc brakes, heat exchangers, gas turbines and chemical reactors, is still their quite high price. Several main processes, including chemical vapor infiltra￾tion (CVI), liquid and solid infiltration and polymer conversion, have been employed for fabrication of this new class of the materials [3–11]. The isothermal CVI method is considered as one of the most promising techniques because its major advan￾tage is the lower thermal and mechanical stress achieved during ∗ Corresponding author. Tel.: +86 24 83978056. E-mail address: jydeng@imr.ac.cn (J. Deng). the densification process, and also because large numbers of varying preforms are easily accommodated in a reactor [3–5]. However, the process has several disadvantages for fabricating the composites: low deposition rate, time-consuming treatment, severe corrosion of vacuum pump, and consequent high cost. The liquid silicon infiltration (LSI) is an attractive method for preparation of the C/SiC composites due to an effective cost [6–8]. The method is infiltration of molten Si into low den￾sity C/C and transformation by a reactive sintering process into the dense materials. Nevertheless, the residual Si metal and the etched fibers will result in poor high-temperature strength of the composites. Another interesting method is the liquid polymer impregnation (LPI) method based on organometallic precur￾sors such as polycarbosilane [9–11]. However, high numbers of impregnation/pyrolysis steps employed for manufacturing a relatively dense material are required because of the high poros￾ity of the initial fiber preform and the low ceramic yield of the precursor. Therefore, it is required to develop a low-cost process for fabricating the high-performance thermo-structural composites. Such an attempt has been made and a novel low-cost manufac￾turing method has been developed, primarily for manufacturing of the C/C, C/SiC and SiC/SiC composites [12,13]. In this paper, large thickness, high density and near-stoichiometric matrix C/SiC composites were fabricated by a simple heaterless 0921-5093/$ – see front matter © 2007 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2007.02.037

S Tang et al. Materials Science and Engineering A 465 (2007)290-294 chemical vapor infiltration(HCVI) technique in a very short ried out by three-point bending, using 80 mm x 10 mm x 6mm time; and the materials were characterized in terms of density, size longitudinal specimens with a span equal to 70 mm and microstructure,composition,mechanical properties and thermal loaded at I mm/min. Compressive tests were performed on 10 mm x 10 mm x 10 mm specimens with a loading speed of I mm/min Fracture toughness was measured on longitudinal 2. Experimental samples of 450mm x 100mm x 4.8 mm at 0.05 mm/min to reduce the deformation rate for stable crack propagation through In the present work, a three-stage and easily controlled pro- a single-edge notch beam(SENB)test using a 40 mm span, a ess was used to fabricate the C/SiC composites. At first, the 5 mm notch depth and a 0. 2 mm notch width[15]. Five, five and 2D preform with a size of 280 mm x 70 mm x 8 mm was fabri- three samples were measured for the flexural, compressive and cated by altematively stacked weftless plies and short-cut-fiber SENB tests, respectively. In-plane CTE from 200 to 1000C was webs using a needle-punching technique and two successive measured on bars with dimensions of 65 mm x8 mm x mm plies were oriented at an angle of 90. Carbon fiber types of them Through-thickness specific heat and TC were obtained through were both PAN-based carbon fiber T700, 12 K tow(from Toray, 12.0 mm x 2.5 mm samples on a FlashlineTM-5000 thermal Japan). The fiber contents of the weftless plies, the webs and the properties analyzer needle fibers were 24.0, 4.5 and 1.5 vol %, respectively. The pre- form was clamped in a graphite clamp with a purpose of wiping 3. Results and discussion off the fiber sizing and maintaining the preform shape. The pro- cessing temperature was increased to 1200C and kept there for Fig. I shows a typical SEM micrograph of the polished cross- 2 h in a vacuum furnace. Secondly, in order to produce a pyro- sections of as-processed C/SiC composites. The cross-section carbon(PyC)interphase and increase the TC, a small amount of area includes O fiber plies, short-cut-fiber webs, 90 fiber plies, Pyc was added to the treated preform by pyrolysing natural gas needle fibers, dense SiC matrix and pores, which is mainly dom- for 10 h at 1000oC in an ICVI apparatus. In the third stage, the inated by well-consolidated parts(Fig. 1(a)). A large amount of treated preforms, clamped between two graphite electrodes for matrix is formed and a few isolated macropores about from 50 to directly heating by passing an electric current in a 50kw cold- 300 um are retained in both the inter-bundles and the inter-ply wall and normal-pressure furnace, were rapidly densified by regions. In a traditional ICVI prepared sample, this is a more forming a SiC matrix by the HCVItechnique. The technique par- typically observed phenomenon and the sizes of these pores are ially overcame the limits of mass transfer and chemical kinetics larger. With a further observation of the intra-bundle, the dense in the traditional ICVI process since a strong temperature gra- inter-fiber matrix can be identified and some residual microp- dient was formed in the fiber preform and an electromagnetic ores are found(Fig. 1(b). From the micrograph, it is possible ield was established around the fiber [14], while one piece of to measure the Sic coating thickness in the inter-fiber pores sample can be prepared one time. Methyltrichlorosilane(MTS, The linear deposition rate in the pores, which is a more practical CH3 SiCl3) was used as a precursor for depositing the SiC matrix, method to evaluate the efficiency of the process, is obtained by and was carried by bubbling hydrogen gas into the chamber. dividing the coating thickness by the infiltration time. The lin- Argon was used as a dilute gas to slow the reaction rate. The ear deposition rates of the SiC matrix within the large inter-fiber leposition conditions were as follows: the hydrogen-to-MTs pores are 0.48-1.01 um/h, partially depending on the spaces temperature below 1000.C and the deposition time 25h. mon mole ratio 1-2, the total flow rate 0.17m/h, the deposition between these pores, which indicates the high deposition effi- ciency. The overall deposition rate can be calculated by dividin Bulk density and open porosity of the C/SiC composites the weight gain per unit time by the volume, and is 0. 10 g/hcm were evaluated by a water immersion technique. Microstruc- As expected, the composites obtain a density of 2.32 g/cm after ture and composition were analyzed using scanning electron a 25-h deposition and its open porosity is only 9.8% microscopy(SEM), energy dispersive spectroscopy (EDS)and The quantitatively chemical compositions of C and Si in the electron probe microanalysis(EPMA). Flexural tests were car- SiC matrix at the center and the edge of the sample according fiber ply Needle~fibe Fig. 1. SEM micrographs of: (a) the macropores between the fiber plies and between the fiber bundles and (b)the micropores between the fibers

S. Tang et al. / Materials Science and Engineering A 465 (2007) 290–294 291 chemical vapor infiltration (HCVI) technique in a very short time; and the materials were characterized in terms of density, microstructure, composition, mechanical properties and thermal properties. 2. Experimental In the present work, a three-stage and easily controlled pro￾cess was used to fabricate the C/SiC composites. At first, the 2D preform with a size of 280 mm × 70 mm × 8 mm was fabri￾cated by alternatively stacked weftless plies and short-cut-fiber webs using a needle-punching technique and two successive plies were oriented at an angle of 90◦. Carbon fiber types of them were both PAN-based carbon fiber T700, 12 K tow (from Toray, Japan). The fiber contents of the weftless plies, the webs and the needle fibers were 24.0, 4.5 and 1.5 vol.%, respectively. The pre￾form was clamped in a graphite clamp with a purpose of wiping off the fiber sizing and maintaining the preform shape. The pro￾cessing temperature was increased to 1200 ◦C and kept there for 2 h in a vacuum furnace. Secondly, in order to produce a pyro￾carbon (PyC) interphase and increase the TC, a small amount of PyC was added to the treated preform by pyrolysing natural gas for 10 h at 1000 ◦C in an ICVI apparatus. In the third stage, the treated preforms, clamped between two graphite electrodes for directly heating by passing an electric current in a 50 kW cold￾wall and normal-pressure furnace, were rapidly densified by forming a SiC matrix by the HCVI technique. The technique par￾tially overcame the limits of mass transfer and chemical kinetics in the traditional ICVI process since a strong temperature gra￾dient was formed in the fiber preform and an electromagnetic field was established around the fiber [14], while one piece of sample can be prepared one time. Methyltrichlorosilane (MTS, CH3SiCl3) was used as a precursor for depositing the SiC matrix, and was carried by bubbling hydrogen gas into the chamber. Argon was used as a dilute gas to slow the reaction rate. The deposition conditions were as follows: the hydrogen-to-MTS mole ratio 1–2, the total flow rate 0.17 m3/h, the deposition temperature below 1000 ◦C and the deposition time 25 h. Bulk density and open porosity of the C/SiC composites were evaluated by a water immersion technique. Microstruc￾ture and composition were analyzed using scanning electron microscopy (SEM), energy dispersive spectroscopy (EDS) and electron probe microanalysis (EPMA). Flexural tests were car￾ried out by three-point bending, using 80 mm × 10 mm × 6 mm size longitudinal specimens with a span equal to 70 mm and loaded at 1 mm/min. Compressive tests were performed on 10 mm × 10 mm × 10 mm specimens with a loading speed of 1 mm/min. Fracture toughness was measured on longitudinal samples of 45.0 mm × 10.0 mm × 4.8 mm at 0.05 mm/min to reduce the deformation rate for stable crack propagation through a single-edge notch beam (SENB) test using a 40 mm span, a 5 mm notch depth and a 0.2 mm notch width [15]. Five, five and three samples were measured for the flexural, compressive and SENB tests, respectively. In-plane CTE from 200 to 1000 ◦C was measured on bars with dimensions of 65 mm × 8 mm × 5 mm. Through-thickness specific heat and TC were obtained through Ø 12.0 mm × 2.5 mm samples on a FlashlineTM-5000 thermal properties analyzer. 3. Results and discussion Fig. 1 shows a typical SEM micrograph of the polished cross￾sections of as-processed C/SiC composites. The cross-section area includes 0◦ fiber plies, short-cut-fiber webs, 90◦ fiber plies, needle fibers, dense SiC matrix and pores, which is mainly dom￾inated by well-consolidated parts (Fig. 1(a)). A large amount of matrix is formed and a few isolated macropores about from 50 to 300m are retained in both the inter-bundles and the inter-ply regions. In a traditional ICVI prepared sample, this is a more typically observed phenomenon and the sizes of these pores are larger. With a further observation of the intra-bundle, the dense inter-fiber matrix can be identified and some residual microp￾ores are found (Fig. 1(b)). From the micrograph, it is possible to measure the SiC coating thickness in the inter-fiber pores. The linear deposition rate in the pores, which is a more practical method to evaluate the efficiency of the process, is obtained by dividing the coating thickness by the infiltration time. The lin￾ear deposition rates of the SiC matrix within the large inter-fiber pores are 0.48–1.01m/h, partially depending on the spaces between these pores, which indicates the high deposition effi- ciency. The overall deposition rate can be calculated by dividing the weight gain per unit time by the volume, and is 0.10 g/h cm3. As expected, the composites obtain a density of 2.32 g/cm3 after a 25-h deposition and its open porosity is only 9.8%. The quantitatively chemical compositions of C and Si in the SiC matrix at the center and the edge of the sample according Fig. 1. SEM micrographs of: (a) the macropores between the fiber plies and between the fiber bundles and (b) the micropores between the fibers.

S Tang et al /Materials Science and Engineering A 465(2007)290-294 Table 1 Weight and mole compositions of C and Si in the SiC matrix at the center and the edge of the cross-section of the C/SiC composites according to EPMA Point Edge C(wt %) Si(wt %) SiC(wt %) C(mol %) Si(mol%) C(wt %) Si(wt %) SiC(wt %) C(mol %) Si(mol %) 9839 52.33 0.11 47.57 52.24 51.87 69. 47.82 52.18 Average 753 48.14 51.86 920 Silicon-Ko Carbon.Ka Chlorine and nitrogen Fig 2 SEM micrograph of SiC growth morphology surrounding the carbon Fig. 3. Load and fracture toughness vs displacement curves for the C/Sic fiber and EDS analysis. to the EPMa are exhibited in Table 1. The deposit is reaching the maximum value for the stress, and then the stress stoichiometric or has a very slight silicon excess (Si to C gradually decreases with the displacement increasing. This indi- about 1.08-1.09)and the compositions are similar at the cates a damage-tolerant fracture behavior, which is also verified and at the edge by micrographs of the fracture surface(Fig 4). The fracture Fig 2 presents the PyC interphase and the Sic matrix sur- surface is jagged. There is obvious subsequent breaking of the rounding the carbon fiber and the eds result. It can be clearly whole fiber bundles because the Sic matrix is firstly deposited seen that a distinct PyCinterphase with a thickness about 1.3 um on the surface of the fibers and then on the surface of the fiber is present on the surface of the fiber and a large amount of bundles. These fiber bundles also have jagged surfaces with the Si accumulates in the PyC, gradually increasing from the fiber fiber pull-out and are not broken in one plane. In some areas, periphery to the PyC periphery. A radical magnetic attraction- many long pull-out fibers are observed, which once more points electric deposition mechanism has been proposed to explain the out the damage tolerance of the developed material. It should be rapid growth rate of the SiC matrix [14. The Si diffusion into emphasized that the majority of the pull-out fibers are smooth the PyC interphase should be attributed to the strong electro- and their diameters are about 7-8 um, proving the crack mainly magnetic affinity of the carbon fibers to the radicals. It should occurring at the interface between the fiber and the PyC. Fur be noted that the interface between the PyC and the fiber is ther, it can be obviously seen that in Fig. 4(c)that many holes smooth and reflect well the fiber shape, possibly revealing the and the Pyc interphases leave behind after the fibers are pulled weak interfacial bonding, whereas the interface between the Pyc out. It is well known that carbon fibers and SiC matrix exhibit and the SiC matrix is tortuous and uneven, probably indicating a siginificantly different CTEs. The CTEs for the fibers are a the strong bonding radial CTE (afr)of 7.0 x 10-6oC-I and a longitudinal CTE The mechanical test results are summarized in Table 2. The (an) of-(0. 1-1. 1)x10-6oC-I, and that of the Sic matrix is curves of load and fracture toughness versus displacement of 4. x 10-C(am)[16, 17]. Since afr >am, the fibers have a the C/SiC composites in the SENB tests are shown in Fig. 3. tendency to radically contract within the SiC matrix upon cool- The composites present a slight deviation from linearity before ing with de-cohesion between the fibers and the PyC interphase Density, porosity and mechanical properties of the C/SiC composites prepared by the HCVi process sity(g/cm) Porosity (%6) Flexural strength(MPa ulus(GPa) Compressive streng gth(MPa)Fracture toughness(MPama) 2.32±004 9.8±0.3 163±8 26.5±2.2 304±55(D276±92(1) 6.5±1.2

292 S. Tang et al. / Materials Science and Engineering A 465 (2007) 290–294 Table 1 Weight and mole compositions of C and Si in the SiC matrix at the center and the edge of the cross-section of the C/SiC composites according to EPMA Point Center Edge C (wt.%) Si (wt.%) SiC (wt.%) C (mol.%) Si (mol.%) C (wt.%) Si (wt.%) SiC (wt.%) C (mol.%) Si (mol.%) 1 27.59 70.80 98.39 47.67 52.33 27.19 70.11 97.30 47.57 52.43 2 27.58 70.54 98.12 47.76 52.24 27.65 69.68 97.33 48.13 51.87 3 27.41 69.92 97.33 47.82 52.18 28.40 69.88 98.28 48.72 51.28 Average 27.53 70.42 97.95 47.75 52.25 28.75 69.89 97.64 48.14 51.86 Fig. 2. SEM micrograph of SiC growth morphology surrounding the carbon fiber and EDS analysis. to the EPMA are exhibited in Table 1. The deposit is near￾stoichiometric or has a very slight silicon excess (Si to C ratio, about 1.08–1.09) and the compositions are similar at the center and at the edge. Fig. 2 presents the PyC interphase and the SiC matrix sur￾rounding the carbon fiber and the EDS result. It can be clearly seen that a distinct PyC interphase with a thickness about 1.3 m is present on the surface of the fiber and a large amount of Si accumulates in the PyC, gradually increasing from the fiber periphery to the PyC periphery. A radical magnetic attraction￾electric deposition mechanism has been proposed to explain the rapid growth rate of the SiC matrix [14]. The Si diffusion into the PyC interphase should be attributed to the strong electro￾magnetic affinity of the carbon fibers to the radicals. It should be noted that the interface between the PyC and the fiber is smooth and reflect well the fiber shape, possibly revealing the weak interfacial bonding, whereas the interface between the PyC and the SiC matrix is tortuous and uneven, probably indicating the strong bonding. The mechanical test results are summarized in Table 2. The curves of load and fracture toughness versus displacement of the C/SiC composites in the SENB tests are shown in Fig. 3. The composites present a slight deviation from linearity before Fig. 3. Load and fracture toughness vs. displacement curves for the C/SiC composites. reaching the maximum value for the stress, and then the stress gradually decreases with the displacement increasing. This indi￾cates a damage-tolerant fracture behavior, which is also verified by micrographs of the fracture surface (Fig. 4). The fracture surface is jagged. There is obvious subsequent breaking of the whole fiber bundles because the SiC matrix is firstly deposited on the surface of the fibers and then on the surface of the fiber bundles. These fiber bundles also have jagged surfaces with the fiber pull-out and are not broken in one plane. In some areas, many long pull-out fibers are observed, which once more points out the damage tolerance of the developed material. It should be emphasized that the majority of the pull-out fibers are smooth and their diameters are about 7–8 m, proving the crack mainly occurring at the interface between the fiber and the PyC. Fur￾ther, it can be obviously seen that in Fig. 4(c) that many holes and the PyC interphases leave behind after the fibers are pulled out. It is well known that carbon fibers and SiC matrix exhibit a siginificantly different CTEs. The CTEs for the fibers are a radial CTE (αfr) of 7.0 × 10−6 ◦C−1 and a longitudinal CTE (αfl) of −(0.1–1.1) × 10−6 ◦C−1, and that of the SiC matrix is 4.8 × 10−6 ◦C−1(αm) [16,17]. Since αfr > αm, the fibers have a tendency to radically contract within the SiC matrix upon cool￾ing with de-cohesion between the fibers and the PyC interphase Table 2 Density, porosity and mechanical properties of the C/SiC composites prepared by the HCVI process Density (g/cm3) Porosity (%) Flexural strength (MPa) Modulus (GPa) Compressive strength (MPa) Fracture toughness (MPa m1/2) 2.32 ± 0.04 9.8 ± 0.3 163 ± 8 26.5 ± 2.2 304 ± 55(||) 276 ± 92(⊥) 6.5 ± 1.2

S Tang et al./ Materials Science and Engineering A 465 (2007)290-294 Fig 4. Fracture surfaces images of:(a)fiber bundles pull-out, (b) long fibers pull-out and (c) residual holes and PyC interphases under a flexural loading 1 10020030040050060070080090010001100 10020030040050060070080090010001100 Fig. 5. (a)CTE-temperature curve and(b)specific heat and TC-temperature curves for the C/SiC composites. yeak interfacial bonding between the fibers and the PyC formed. from 847 J/kgC at 200C to 1292 J/kgC at 1000C. The Moreover, almost no Si diffusion into the fiber also contributes of the C/SiC composites is relatively low; and it decreases with to the weak bonding. However, the interfacial bonding between temperature increasing before 500"C and then slightly increases be stronger since a sig- from 500 to 1000 C. The characteristic is different for ceramics nificant amount of Si has diffused into the Pyc interphase, as and carbon, whose TC decreases with increasing temperature. It shown in Fig. 2. So the fibers are more easily pulled out at the has been well known that tC of a well-crystallized Sic matrix interface between the fibers and the PyC than at the interface is expected to be higher than that of a poor-crystallized or amor- between the PyC and the matrix with the result that an advanc- phous SiC matrix. So the Sic matrix prepared by the HCVI ing longitudinal matrix crack can propagate and deflect in the process should possess a high TC since it possesses a well- PyC interphases, and then develop along the fibers, which may crystallized Sic matrix and a pronounced anisotropy in the improve the damage-tolerant behavior. (11 1)direction [14]. However, the TC also depends on the fiber The in-plane CtE of the C/SiC composites in the range architecture. A small amount of fiber content and the large inter 200-1000C is exhibited in Fig. 5(a). The measured CTE, layer pores can contribute to the low TC in the through-thickness (1. 14-1.93)x100C-, is significantly lower than silicon car direction since the fiber architecture is fabricated by the weft- bide monolithic ceramic, 4.8x 10-6oC-I. The slope of the CTE less plies with 12 K large tows and the super-thin webs using in the range 200-500C is obviously lower than that in the a needle technique. The small increase of the TC after 500C range 500-1000C. Compressibility of the fibers and stretch may be attributed to the continuous heat transmission of the SiC of the matrix can be produced during cooling from the process- matrix due to the cracks sealing at 500C ing temperature owing to their different CTEs. The compressive stress in the fibers is released as the matrix expands, and hence 4. Conclusions y the expa Since the expansion of the matrix may be partially counteracted The C/SiC composites have been rapidly produced by a HCVI by the thermal stress cracks sealing in a low temperature, the process. The inter-layer and inter-bundle pores are well den slope of the CTE of the composites is lower at 200-500C than sified and a small porosity is obtained. The composites show that at 500-1000C. From the above result, it may be inferred a damage-tolerant fracture behavior and the cracking mainly that the crack healing temperature is about 500C. Fig. 5(b) occurs at the interface between the fibers and the Pyc due presents the curves of specific heat and TC-temperature in the to their different CTE and many Si-atoms diffused into the range 200-1000C. The specific heat of the C/SiC composites PyC interphase. Its flexural strength, longitudinal/transverse

S. Tang et al. / Materials Science and Engineering A 465 (2007) 290–294 293 Fig. 4. Fracture surfaces images of: (a) fiber bundles pull-out, (b) long fibers pull-out and (c) residual holes and PyC interphases under a flexural loading. Fig. 5. (a) CTE-temperature curve and (b) specific heat and TC-temperature curves for the C/SiC composites. due to a large amount of Si diffusion into the interphase, hence a weak interfacial bonding between the fibers and the PyC formed. Moreover, almost no Si diffusion into the fiber also contributes to the weak bonding. However, the interfacial bonding between the SiC matrix and the PyC should be stronger since a sig￾nificant amount of Si has diffused into the PyC interphase, as shown in Fig. 2. So the fibers are more easily pulled out at the interface between the fibers and the PyC than at the interface between the PyC and the matrix with the result that an advanc￾ing longitudinal matrix crack can propagate and deflect in the PyC interphases, and then develop along the fibers, which may improve the damage-tolerant behavior. The in-plane CTE of the C/SiC composites in the range 200–1000 ◦C is exhibited in Fig. 5(a). The measured CTE, (1.14–1.93) × 10−6 ◦C−1, is significantly lower than silicon car￾bide monolithic ceramic, 4.8 × 10−6 ◦C−1. The slope of the CTE in the range 200–500 ◦C is obviously lower than that in the range 500–1000 ◦C. Compressibility of the fibers and stretch of the matrix can be produced during cooling from the process￾ing temperature owing to their different CTEs. The compressive stress in the fibers is released as the matrix expands, and hence the in-plane CTE is controlled by the expansion of the matrix. Since the expansion of the matrix may be partially counteracted by the thermal stress cracks sealing in a low temperature, the slope of the CTE of the composites is lower at 200–500 ◦C than that at 500–1000 ◦C. From the above result, it may be inferred that the crack healing temperature is about 500 ◦C. Fig. 5(b) presents the curves of specific heat and TC-temperature in the range 200–1000 ◦C. The specific heat of the C/SiC composites increases by about a factor of 1.5 over this temperature range, from 847 J/kg ◦C at 200 ◦C to 1292 J/kg ◦C at 1000 ◦C. The TC of the C/SiC composites is relatively low; and it decreases with temperature increasing before 500 ◦C and then slightly increases from 500 to 1000 ◦C. The characteristic is different for ceramics and carbon, whose TC decreases with increasing temperature. It has been well known that TC of a well-crystallized SiC matrix is expected to be higher than that of a poor-crystallized or amor￾phous SiC matrix. So the SiC matrix prepared by the HCVI process should possess a high TC since it possesses a well￾crystallized SiC matrix and a pronounced anisotropy in the 111 direction [14]. However, the TC also depends on the fiber architecture. A small amount of fiber content and the large inter￾layer pores can contribute to the low TC in the through-thickness direction since the fiber architecture is fabricated by the weft￾less plies with 12 K large tows and the super-thin webs using a needle technique. The small increase of the TC after 500 ◦C may be attributed to the continuous heat transmission of the SiC matrix due to the cracks sealing at 500 ◦C. 4. Conclusions The C/SiC composites have been rapidly produced by a HCVI process. The inter-layer and inter-bundle pores are well den￾sified and a small porosity is obtained. The composites show a damage-tolerant fracture behavior and the cracking mainly occurs at the interface between the fibers and the PyC due to their different CTE and many Si-atoms diffused into the PyC interphase. Its flexural strength, longitudinal/transverse

S Tang et al /Materials Science and Engineering A 465(2007)290-294 compressive strength, fracture toughness, and CTE and TC in the [5]JQMa, Y.D. Xu,LT.Zhang, LFCheng,JJNie, HLi,Mater.Lett.61 range 200-1000oC are 163 MPa, 304/276 MPa, 6.5 MPam/2 (2007)312-315 (1.09-1.93)x10-0K-,8.34-6.56 W/m K, respectively. The [6] FH Gem, R. Kochendorfer, Compos. Part A 28(1997)335-364 sealing temperature of the thermal stress cracks is inferred to be [7 w Krenkel, Int JAppl.Ceram. Tech 1(2004)188-200 500C because of the different slopes of the Cte and TC curves (9) M.Z. Berbon, D R. Dietrich, D B. Marshall, Am Ceram Soc. 84(2001) in the 200-500 and 500-1000oC regions From the technolog 2229-2234 ical point of view, the fabrication process offers the advantage [10] G.B. Zheng, H Sano, Y. Uchiyama, K. Kobayashi, H.M. Cheng, JMater. Sci.34(1999)827-834. of low-cost and short-time production of the C/SiC composites [11]MFGonon, S.Hampshire,J.Eur.Ceram Soc.19(1999)285-291 with good properties [12] W.C. Liu, J.Y. Deng, H F Du, F. Liu, China Patent No 99 1 22649.6(2003) [13] S.F. Tang J.Y. Deng, H F Du, w.C. Liu, K Yang, J Am Ceram Soc. References (2005)3253-3257 [14]SF. Tang, J.Y. Deng, H.F. Du, w.C. Liu, in pre [1 KJ. Torben, B. Povl, J Am Ceram Soc. 84(2001)1043-1051 [15] Z.D. Guan, Z.T. Zhang, J.S. Jiao, Physical Properties of Inorganic Materi- t.R. Naslain, Compos. Sci. Technol. 64(2004)155-170. als, Tsinghua University Press, Beijing. 1992. pp 57-67 Mall, J.M. Engesser, Compos. Sci. Technol. 66(2006)863-874 [16] Y Xu, L Zhang, L Cheng, D. Yan, Carbon 36(1998)1051-1056. [4]R.R. Naslain, R. Pailler, X. Bourrat, S. Bertrand, F. Heurtevent, P. Dupel, [17] F. Lamouroux, X. Bourrat, R. Naslain, J. Sevely, Carbon 31(1993) F Lamouroux, Solid State ionics 141-142(2001)541-548 1273-128

294 S. Tang et al. / Materials Science and Engineering A 465 (2007) 290–294 compressive strength, fracture toughness, and CTE and TC in the range 200–1000 ◦C are 163 MPa, 304/276 MPa, 6.5 MPa m1/2, (1.09–1.93) × 10−6 K−1, 8.34–6.56 W/m K, respectively. The sealing temperature of the thermal stress cracks is inferred to be 500 ◦C because of the different slopes of the CTE and TC curves in the 200–500 and 500–1000 ◦C regions. From the technolog￾ical point of view, the fabrication process offers the advantage of low-cost and short-time production of the C/SiC composites with good properties. References [1] K.J. Torben, B. Povl, J. Am. Ceram. Soc. 84 (2001) 1043–1051. [2] R.R. Naslain, Compos. Sci. Technol. 64 (2004) 155–170. [3] S. Mall, J.M. Engesser, Compos. Sci. Technol. 66 (2006) 863–874. [4] R.R. Naslain, R. Pailler, X. Bourrat, S. Bertrand, F. Heurtevent, P. Dupel, F. Lamouroux, Solid State Ionics 141–142 (2001) 541–548. [5] J.Q. Ma, Y.D. Xu, L.T. Zhang, L.F. Cheng, J.J. Nie, H. Li, Mater. Lett. 61 (2007) 312–315. [6] F.H. Gern, R. Kochendorfer, Compos.: Part A 28 (1997) 335–364. [7] W. Krenkel, Int. J. Appl. Ceram. Tech. 1 (2004) 188–200. [8] W. Krenkel, Ceram. Eng. Sci. P 22 (2001) 443–454. [9] M.Z. Berbon, D.R. Dietrich, D.B. Marshall, J. Am. Ceram. Soc. 84 (2001) 2229–2234. [10] G.B. Zheng, H. Sano, Y. Uchiyama, K. Kobayashi, H.M. Cheng, J. Mater. Sci. 34 (1999) 827–834. [11] M.F. Gonon, S. Hampshire, J. Eur. Ceram. Soc. 19 (1999) 285–291. [12] W.C. Liu, J.Y. Deng, H.F. Du, F. Liu, China Patent No. 99 1 22649.6 (2003). [13] S.F. Tang, J.Y. Deng, H.F. Du, W.C. Liu, K. Yang, J. Am. Ceram. Soc. 88 (2005) 3253–3257. [14] S.F. Tang, J.Y. Deng, H.F. Du, W.C. Liu, in preparation. [15] Z.D. Guan, Z.T. Zhang, J.S. Jiao, Physical Properties of Inorganic Materi￾als, Tsinghua University Press, Beijing, 1992, pp. 57–67. [16] Y. Xu, L. Zhang, L. Cheng, D. Yan, Carbon 36 (1998) 1051–1056. [17] F. Lamouroux, X. Bourrat, R. Naslain, J. Sevely, Carbon 31 (1993) 1273–1288