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/transverseS. 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