B. Wilshire, M.R. Bache/ Journal of the European Ceramic Sociery 27(2007)4603-4611 fied as the total creep strain to failure(Ef) Continuously decaying creep strain/time curves of the form shown in Figs. I and 7 therefore terminate when t=tf ande 2 Em when a=Ef. In this context, with the SiCr-SiC composite, the creep duc tilities are very low(Fig. 6)so the creep curves terminate early (Fig. 7), giving high Em and low tr values(Figs. 3 and 4) Although the initial rates of creep strain accumulation are similar because the fibre reinforcement is essentially the same, the creep ductilities of the SiCf-SiBC material are higher than those for the Cr-SiC samples(Fig. 6). Creep therefore continues until the larger Er values are attained, so lower creeprates and much longer creep lives are displayed under the same test conditions( Fig. 7) Then, by replacing the NicalonTM NLM 202 with stronger Hi- 1 mm Nicalon fibres, the rates of creep strain accumulation are also decreased(Fig. 7), so the large Ef values(Fig. 6) lead to the creep and creep rupture strength of the HNSiCr-SiBC compos- (b) ite being considerably superior to the properties displayed by the Sicr-SiBC specimens(Figs. 3 and 4). Yet, while substan- tial performance gains would be anticipated by incorporation of stronger fibres the matrices contribute little to the stress- bearing capabilities of these CFCMCs Even so, with nominally identical fibre reinforcement, replacing SiC with either SiBC or Al2 O3 matrices leads to significant strength enhancements (Figs. 3 and 4). This observation can then be interpreted by con- sidering the deformation and damage processes governing strain accumulation and failure 3.4. Creep deformation and damage processes 100μm On applying a tensile load to a'textile'CFCMC, the inter- Fig 8. Scanning electron micrographs showing crack development through the woven longitudinal fibre bundles extend and straighten in the transverse(90 )fibre bundles, by-passing the fibres as the cracks grow through stress direction However as with the individual fibres. b the (a) the alumina matrix of the HNSiCr-Al2O3 and(b) the SiBC matrix of the creep strengths of fibre bundles also vary. 8 Hence, the weak- HNSiCr-SiBC composites In both cases, the tensile stress axis is vertical est fibre regions deform most easily, so the high initial creep rates decrease with time as load is transferred progressively to fracture takes place by fibre pull-out(Fig. 9b). In this way, stronger bundles. This deformation is accompanied by crack for- the longitudinal 0 fibres control the rates of strain accumu- crack development has little effect on the overall strength of the oxidation-assisted fibre failure and creep ductility. 8 rates c mation in the brittle matrices but, because the matrices are weak, lation and crack growth, while the matrices affect the rates of composite. Consequently, the creep rate continues to decrease With the SiC-SiC samples,6 the dominant macro-crack with time(Figs. I and 7), with the rate of strain accumulation nucleates at surface macro-pores, with direct oxygen penetra decreasing as the stress and temperature decrease. tion along the crack leading to low ductility failure(Fig. 7). In As the longitudinal bundles extend and straighten, the result- contrast, with the SiCr-SiBC material, glass formation limits ing complex stress state leads to crack formation within and oxidation-assisted fibre failure, giving higher creep ductilities between the 00 and 90 fibre tows, although composite failure is (Fig. 7). Further performance gains are then achieved with governed by the growth of"'macro-cracks'along planes normal the HNSiCr-SiBC product, when the benefits of the enhanced to the tensile axis. 8 These'tunnellingcracks easily by-pass SiBC matrices are combined with the reduced rates of strain fibres in the 90 bundles(Fig. 8)but, on penetrating into the accumulation achieved by replacing NicalonM NLM202 with 0 tows, the cracks become bridged by unbroken longitudinal Hi-Nicalon M fibres(Fig. 7) fibres. Cracks can then open and extend only at rates determined With alumina matrices. residual stress-induced micro-cracks by the creep resistance of the bridging fibres, accounting for the are present in the as-produced samples, so many small cracks dependence of tr on Em(Fig. 5) develop throughout the gauge length of the AlO3-matrix com- Unfortunately, in oxidizing environments, matrix cracking posites as creep proceeds. Indirect oxygen ingress through promotes oxygen ingress, causing premature failure of the crack- the micro-cracked matrices, coupled with the more oxidation bridging fibres and accelerating crack growth. 8 In general, the resistant double BN/SiC interfaces, then results in a resistance dominant crack causing failure is surface nucleated(Fig. 9a), to oxidation-assisted fibre failure equivalent to that of sibC developing until the stress on the remaining unbroken cross- matrices. Hence, similar creep and creep rupture strengths section of the composite reaches the critical level at which are exhibited by the Sicr-AlzO3 and Sicr-SiBC materials, asB. Wilshire, M.R. Bache / Journal of the European Ceramic Society 27 (2007) 4603–4611 4607 fied as the total creep strain to failure (εf). Continuously decaying creep strain/time curves of the form shown in Figs. 1 and 7 therefore terminate when t = tf and ε˙ ∼= ε˙m when ε = εf. In this context, with the SiCf–SiC composite, the creep duc￾tilities are very low (Fig. 6) so the creep curves terminate early (Fig. 7), giving high ε˙m and low tf values (Figs. 3 and 4). Although the initial rates of creep strain accumulation are similar because the fibre reinforcement is essentially the same, the creep ductilities of the SiCf–SiBC material are higher than those for the SiCf–SiC samples (Fig. 6). Creep therefore continues until the larger εf values are attained, so lower creep rates and much longer creep lives are displayed under the same test conditions (Fig. 7). Then, by replacing the NicalonTM NLM 202 with stronger Hi￾NicalonTM fibres, the rates of creep strain accumulation are also decreased (Fig. 7), so the large εf values (Fig. 6) lead to the creep and creep rupture strength of the HNSiCf–SiBC compos￾ite being considerably superior to the properties displayed by the SiCf–SiBC specimens (Figs. 3 and 4). Yet, while substan￾tial performance gains would be anticipated by incorporation of stronger fibres, the matrices contribute little to the stress￾bearing capabilities of these CFCMCs. Even so, with nominally identical fibre reinforcement, replacing SiC with either SiBC or Al2O3 matrices leads to significant strength enhancements (Figs. 3 and 4). This observation can then be interpreted by con￾sidering the deformation and damage processes governing strain accumulation and failure. 3.4. Creep deformation and damage processes On applying a tensile load to a ‘textile’ CFCMC, the inter￾woven longitudinal fibre bundles extend and straighten in the stress direction. However, as with the individual fibres,16 the creep strengths of fibre bundles also vary.18 Hence, the weak￾est fibre regions deform most easily, so the high initial creep rates decrease with time as load is transferred progressively to stronger bundles. This deformation is accompanied by crack for￾mation in the brittle matrices but, because the matrices are weak, crack development has little effect on the overall strength of the composite. Consequently, the creep rate continues to decrease with time (Figs. 1 and 7), with the rate of strain accumulation decreasing as the stress and temperature decrease. As the longitudinal bundles extend and straighten, the result￾ing complex stress state leads to crack formation within and between the 0◦ and 90◦ fibre tows, although composite failure is governed by the growth of ‘macro-cracks’ along planes normal to the tensile axis.8 These ‘tunnelling’ cracks easily by-pass fibres in the 90◦ bundles (Fig. 8) but, on penetrating into the 0◦ tows, the cracks become bridged by unbroken longitudinal fibres. Cracks can then open and extend only at rates determined by the creep resistance of the bridging fibres, accounting for the dependence of tf on ε˙m (Fig. 5). Unfortunately, in oxidizing environments, matrix cracking promotes oxygen ingress, causing premature failure of the crack￾bridging fibres and accelerating crack growth.8 In general, the dominant crack causing failure is surface nucleated (Fig. 9a), developing until the stress on the remaining unbroken cross￾section of the composite reaches the critical level at which Fig. 8. Scanning electron micrographs showing crack development through the transverse (90◦) fibre bundles, by-passing the fibres as the cracks grow through: (a) the alumina matrix of the HNSiCf–Al2O3 and (b) the SiBC matrix of the HNSiCf–SiBC composites. In both cases, the tensile stress axis is vertical. fracture takes place by fibre pull-out (Fig. 9b). In this way, the longitudinal 0◦ fibres control the rates of strain accumu￾lation and crack growth, while the matrices affect the rates of oxidation-assisted fibre failure and creep ductility.8 With the SiCf–SiC samples,6 the dominant macro-crack nucleates at surface macro-pores, with direct oxygen penetra￾tion along the crack leading to low ductility failure (Fig. 7). In contrast, with the SiCf–SiBC material, glass formation limits oxidation-assisted fibre failure, giving higher creep ductilities (Fig. 7). Further performance gains are then achieved with the HNSiCf–SiBC product, when the benefits of the enhanced SiBC matrices are combined with the reduced rates of strain accumulation achieved by replacing NicalonTM NLM202 with Hi-NicalonTM fibres (Fig. 7). With alumina matrices, residual stress-induced micro-cracks are present in the as-produced samples,19 so many small cracks develop throughout the gauge length of the Al2O3–matrix com￾posites as creep proceeds.7 Indirect oxygen ingress through the micro-cracked matrices, coupled with the more oxidation￾resistant double BN/SiC interfaces, then results in a resistance to oxidation-assisted fibre failure equivalent to that of SiBC matrices. Hence, similar creep and creep rupture strengths are exhibited by the SiCf–Al2O3 and SiCf–SiBC materials, as