B. Wilshire, F. Carrero/ Materials Science and Engineering 4272 (1999)38-44 With the SiCw/Al,O3 ceramics, as with the present the creep strain/ time plots become less pronounced as SiCr AlO, composite, the decrease in n value with the creep temperature is raised(Fig. 2), the im' t e ratio decreasing applied stress can be explained in terms of increases as the primary strain (Ep) decreases with in- microstructural instability rather than mechanism creasing test temperature(Fig. 10) changes. Oxidation of the sic whiskers can increase the The fact that a increases with decreasing stress(Fig quantities of intergranular viscous phases present, 11)can then be explained on the basis that, for both the which would reduce the creep strength with increasing whisker and fibre reinforced CMCs, fracture occurs by est duration, especially at high creep temperatures. the development of a crack long enough to propagate Grain boundary and interfacial amorphous phases are rapidly under the prevailing test conditions, with the detrimental to creep resistance [5, 7], accounting for the critical crack length increasing with decreasing applied creep strength of SICw/Al2O3 materials being lower in stress. At high stresses, when the critical crack lengths air than in inert atmospheres [10]. This sensitivity to are relatively short, low ductility failures are recorded oxidation is then consistent with the observation that. In contrast formation of the critical crack involves for the 25 vol. whisker reinforced alumina tested in progressive growth and link up of shorter cracks as the tension, the alumina-rich oxidized regions at the speci test duration increases. Since cracks develop more uni- men surfaces increased in thickness with increasing test formly throughout the testpieces at lower stresses and duration and increasing test temperature [12]. On this temperatures, creep must continue for longer times to basis, in line with the results obtained for the present reach higher strains in order to create a crack long SiCHAl2O3 material, the information reported for SiCwl enough to cause failure by fibre pull-out, so that er and Al2O3 composites supports the view [27] that compari- Em'tr increase with decreasing stress at each creep tem- sons of measured and theoretical n and @c values perature(Figs. 10 and ll) cannot provide an unambiguous indication of the dom For the SiCr/Al2O3 composite, while tr is a function inant creep mechanism. of im(Eq (3)), the gradients of the log im/log tr relation Although traditional methods for identifying creep ships are approximately 0.75(Fig 9). Consequently, the mechanisms are unsatisfactory, the data comparisons in n values derived from the stress/creep rate plots(Fig 3) Fig. 5 show that creep of the 0/90 SiC Al2O3 com- are lower than the equivalent stress dependences calcu posite is controlled by the fibres. In turn, creep of the lated from the stress rupture data(Fig 8). For example fibres results in matrix cracking, with the longer cracks at the highest stress levels imposed at 1473, 1573 and developing normal to the tensile axis being bridged by 1673 K, the maximum n values are around 50 for the longitudinal fibres. While oxygen penetration promotes log a/log Em plots in Fig. 3, compared with the equiva failure of the crack-bridging fibres, the dependence of lent maximum values of about 65 for the log a/log tr the rupture life on the minimum creep rate(Eq.(3)) data in Fig 8. Even so, especially at the higher creep confirms that the rate of crack growth and the time to temperatures, both the log a/log Em and log a/log tr fracture are determined by the rate of creep strain plots curve such that the large stress exponents and accumulation. Clearly, as the crack-bridging fibres fail, activation energies found in short-duration tests de- he numbers of fibres supporting the creep loads de- crease with decreasing stress. For this reason, linea crease with increasing creep strain. However, the local- extrapolation of results obtained at high stresses overes- ized increases in stress due to crack development are timates low-stress performance, emphasizing the engi not sufficient to give a discernible acceleration in creep neering design requirement for reliable determination of rate before fracture occurs. Consequently, failure takes long-term property values place by the development of a crack long enough to Other information may be derived from the curving cause fracture by fibre pull-out(Fig. 7), terminating the log o/log Em and log a/log tr plots. For example, since decaying primary curve with no clearly-defined tertiary the n value usually increases with increasing stress, the age(Fig. 2) maximum recorded n value must depend on the stress- While tertiary stage tually absent with both bearing capabilities of the material. Thus, for polycrys- the fibre and whisker orced alumina-matrix ceram- talline magnesia, the maximum measured n values cs, distinct primary stages are found with the Sic increased from approximately 2 to approximately 7 as AlO3 composite(Fig. 2) but not with the SiCwAlO3 the strength and ductility were improved by modifying material [12]. Even so, for both types of creep curve, the fabrication routes[28]. On this basis, the attainment the total creep strain to failure can be expressed as of n values of 50 or more at high stresses(Figs. 3 and f=印+-mr 8)emphasizes the impressive creep and creep fracture resistance of the present 0/90% SiCH/Al2O, composite where ap is the primary creep strain. Thus, with the high temperatures. whisker-reinforced alumina, sp=0 so that = Em'tr ractical viewpoint, the creep and creep (Fig. 10). With the fibre-reinforced alumina, Em'tr also fracture strengths of the fibre-reinforced alumina are increases with increasing ar but, since the curvatures of markedly superior to those of whisker-reinforced prodB. Wilshire, F. Carren˜o / Materials Science and Engineering A272 (1999) 38–44 43 With the SiCw/Al2O3 ceramics, as with the present SiCf /Al2O3 composite, the decrease in n value with decreasing applied stress can be explained in terms of microstructural instability rather than mechanism changes. Oxidation of the SiC whiskers can increase the quantities of intergranular viscous phases present, which would reduce the creep strength with increasing test duration, especially at high creep temperatures. Grain boundary and interfacial amorphous phases are detrimental to creep resistance [5,7], accounting for the creep strength of SiCw/Al2O3 materials being lower in air than in inert atmospheres [10]. This sensitivity to oxidation is then consistent with the observation that, for the 25 vol.% whisker reinforced alumina tested in tension, the alumina-rich oxidized regions at the speci￾men surfaces increased in thickness with increasing test duration and increasing test temperature [12]. On this basis, in line with the results obtained for the present SiCf /Al2O3 material, the information reported for SiCw/ Al2O3 composites supports the view [27] that compari￾sons of measured and theoretical n and Qc values cannot provide an unambiguous indication of the dom￾inant creep mechanism. Although traditional methods for identifying creep mechanisms are unsatisfactory, the data comparisons in Fig. 5 show that creep of the 0/90° SiCf /Al2O3 com￾posite is controlled by the fibres. In turn, creep of the fibres results in matrix cracking, with the longer cracks developing normal to the tensile axis being bridged by longitudinal fibres. While oxygen penetration promotes failure of the crack-bridging fibres, the dependence of the rupture life on the minimum creep rate (Eq. (3)) confirms that the rate of crack growth and the time to fracture are determined by the rate of creep strain accumulation. Clearly, as the crack-bridging fibres fail, the numbers of fibres supporting the creep loads de￾crease with increasing creep strain. However, the local￾ized increases in stress due to crack development are not sufficient to give a discernible acceleration in creep rate before fracture occurs. Consequently, failure takes place by the development of a crack long enough to cause fracture by fibre pull-out (Fig. 7), terminating the decaying primary curve with no clearly-defined tertiary stage (Fig. 2). While tertiary stages are virtually absent with both the fibre and whisker reinforced alumina-matrix ceram￾ics, distinct primary stages are found with the SiCf / Al2O3 composite (Fig. 2) but not with the SiCw/Al2O3 material [12]. Even so, for both types of creep curve, the total creep strain to failure can be expressed as: of=op+o; m·tf (4) where op is the primary creep strain. Thus, with the whisker-reinforced alumina, op$0 so that of$o; m·tf (Fig. 10). With the fibre-reinforced alumina, o; m·tf also increases with increasing of but, since the curvatures of the creep strain/time plots become less pronounced as the creep temperature is raised (Fig. 2), the o; m·tf /of ratio increases as the primary strain (op) decreases with in￾creasing test temperature (Fig. 10). The fact that of increases with decreasing stress (Fig. 11) can then be explained on the basis that, for both the whisker and fibre reinforced CMCs, fracture occurs by the development of a crack long enough to propagate rapidly under the prevailing test conditions, with the critical crack length increasing with decreasing applied stress. At high stresses, when the critical crack lengths are relatively short, low ductility failures are recorded. In contrast, formation of the critical crack involves progressive growth and link up of shorter cracks as the test duration increases. Since cracks develop more uni￾formly throughout the testpieces at lower stresses and temperatures, creep must continue for longer times to reach higher strains in order to create a crack long enough to cause failure by fibre pull-out, so that of and o; m·tf increase with decreasing stress at each creep tem￾perature (Figs. 10 and 11). For the SiCf /Al2O3 composite, while tf is a function of o; m (Eq. (3)), the gradients of the log o; m/log tf relation￾ships are approximately 0.75 (Fig. 9). Consequently, the n values derived from the stress/creep rate plots (Fig. 3) are lower than the equivalent stress dependences calcu￾lated from the stress rupture data (Fig. 8). For example, at the highest stress levels imposed at 1473, 1573 and 1673 K, the maximum n values are around 50 for the log s/log o; m plots in Fig. 3, compared with the equiva￾lent maximum values of about 65 for the log s/log tf data in Fig. 8. Even so, especially at the higher creep temperatures, both the log s/log o; m and log s/log tf plots curve such that the large stress exponents and activation energies found in short-duration tests de￾crease with decreasing stress. For this reason, linear extrapolation of results obtained at high stresses overes￾timates low-stress performance, emphasizing the engi￾neering design requirement for reliable determination of long-term property values. Other information may be derived from the curving log s/log o; m and log s/log tf plots. For example, since the n value usually increases with increasing stress, the maximum recorded n value must depend on the stress￾bearing capabilities of the material. Thus, for polycrys￾talline magnesia, the maximum measured n values increased from approximately 2 to approximately 7 as the strength and ductility were improved by modifying the fabrication routes [28]. On this basis, the attainment of n values of 50 or more at high stresses (Figs. 3 and 8) emphasizes the impressive creep and creep fracture resistance of the present 0/90° SiCf /Al2O3 composite at high temperatures. From a practical viewpoint, the creep and creep fracture strengths of the fibre-reinforced alumina are markedly superior to those of whisker-reinforced prod-