l884 Journal of the American Ceramic Society'Kerans et al. Vol 81. No. 7 8000 8000 7000 7000 6000 6000 乏5000 5000 4000 4000 3000 000 P90/2.2/28 2000 P:50/50/282000 1000 elastic 1000 8.5605 Displacement(micrometers) Fig 3. Push-out curves according to the roughness interface model, with parameters chosen to fit the experimental curve of Fig. 1, assuming that (i) initial cracking is governed by the roughness parameters h =50 nm, d= 50 nm, and G= 28 J/m-and(ii) subsequent cracking is governed by h =50 nm, d= 2. 2 um, and G= 28 J/m", the transition between the two conditions is indicated by the arrowed line In the rough-interface analysis, the interfacial friction will ported composites that show crack deflection along fibers and increase very quickly to extremely high values in the primary good properties. This, in turn, suggests that essentially all other short-period crack regime, then decrease to much-lower values ood"'conventional NicalonTM-reinforced composites in the secondary crack regime. For example, the effective av- debond in a manner similar to that of the untreated-fiber com- erage fiber stress to overcome friction at crack initiation in Fig posites, i.e., along an interface or in a weak layer very near an 4 is-200 MPa, whereas at the Region II/lll transition of the interface, as opposed to through the coating layer itself. rimary crack, this value is -750 MPa. Upon transition to the Perhaps the most important outcome of work related to the long-period crack, it decreases to -50 MPa, then climbs to 85 treated-fiber composites is the suggestion that the fracture en- APa at the Region il/ili transition of the ergy for debonding can be higher than often assumed and still correspond to this model, such as matrix-crack spacing or ten- suggests that the window of properties that can be possessed sile hysteresis loops, will return very high friction values. On by an oxide alternative coating, for example, is larger than such as fiber pushout(in a conventional analysis), will return because the very low fracture energies previously thought to be much-lower friction values necessary will likely be difficult to obtain with oxides or other The differences in the behavior of the treated-fiber compos- alternative materials Ite system considered here, in comparison to a more conven- o It is important to note that the crack-deflection criterion is different interfacial cracking behavior. 2,3 The results of this Hutchinson criterion, for example, was derived for deflection facial crack in untreated-fiber composites has been reported to figure of merit for the competition between crack deflection in be confined to the C/SiO, interface near the fiber surface or the the interface and crack propagation into the second material carbon layers very near to the interface 3, 11 Such composites was the ratio of interfacial (principally mode II)fracture energy have been observed, when measured, to have interfacial frac- to its mode I fracture energy. 5 However, if the relevant issue ture energies of no more than a few joules per square meter. is crack deflection within the coating, the analogous figure of The fiber treatment has been inferred to strengthen the interface merit will be the ratio of the coating(principally mode ID) egion to a level that is above the strength of the pyrocarbon fracture energy(,2)to its mode I fracture energy (T2,2). In coating itself, thereby shifting the fracture to the pyrocarbon, principle, the basic suitability of a material that is intended to which is the next-weakest link 3, 12 The measurements and in promote crack deflection via cleavage can be evaluated inde- terpretation of this work imply that the fracture energy and the pendent of the fiber fracture properties friction both are very much higher than in the untreated-fiber There are interesting implications of the fracture behavior composites and, moreover, much higher than in any other bserved in the materials of the preceding section. Crack de-In the rough-interface analysis, the interfacial friction will increase very quickly to extremely high values in the primary, short-period crack regime, then decrease to much-lower values in the secondary crack regime. For example, the effective av￾erage fiber stress to overcome friction at crack initiation in Fig. 4 is ∼200 MPa, whereas at the Region II/III transition of the primary crack, this value is ∼750 MPa. Upon transition to the long-period crack, it decreases to ∼50 MPa, then climbs to 85 MPa at the Region II/III transition of the long-period crack. Tests that probe the primary-crack regime of specimens that correspond to this model, such as matrix-crack spacing or ten￾sile hysteresis loops, will return very high friction values. On the other hand, tests that probe the secondary-crack regime, such as fiber pushout (in a conventional analysis), will return much-lower friction values. The differences in the behavior of the treated-fiber compos￾ite system considered here, in comparison to a more conven￾tional untreated-fiber composite, have been attributed to the different interfacial cracking behavior.2,3 The results of this work are consistent with that scenario. Specifically, the inter￾facial crack in untreated-fiber composites has been reported to be confined to the C/SiO2 interface near the fiber surface or the carbon layers very near to the interface.3,11 Such composites have been observed, when measured, to have interfacial frac￾ture energies of no more than a few joules per square meter. The fiber treatment has been inferred to strengthen the interface region to a level that is above the strength of the pyrocarbon coating itself, thereby shifting the fracture to the pyrocarbon, which is the next-weakest link.3,12 The measurements and in￾terpretation of this work imply that the fracture energy and the friction both are very much higher than in the untreated-fiber composites and, moreover, much higher than in any other re￾ported composites that show crack deflection along fibers and good properties. This, in turn, suggests that essentially all other ‘‘good’’ conventional Nicalon™-reinforced composites debond in a manner similar to that of the untreated-fiber com￾posites, i.e., along an interface or in a weak layer very near an interface, as opposed to through the coating layer itself. Perhaps the most important outcome of work related to the treated-fiber composites is the suggestion that the fracture en￾ergy for debonding can be higher than often assumed and still perform the necessary function of crack deflection. This result suggests that the window of properties that can be possessed, by an oxide alternative coating, for example, is larger than previously thought. This observation is an important result, because the very low fracture energies previously thought to be necessary will likely be difficult to obtain with oxides or other alternative materials. It is important to note that the crack-deflection criterion is often discussed in a potentially misleading way. The He and Hutchinson criterion, for example, was derived for deflection in a true interface between two materials, and the relevant figure of merit for the competition between crack deflection in the interface and crack propagation into the second material was the ratio of interfacial (principally mode II) fracture energy to its mode I fracture energy.25 However, if the relevant issue is crack deflection within the coating, the analogous figure of merit will be the ratio of the coating (principally mode II) fracture energy (c Gr,z) to its mode I fracture energy (c Gz,z). In principle, the basic suitability of a material that is intended to promote crack deflection via cleavage can be evaluated inde￾pendent of the fiber fracture properties. There are interesting implications of the fracture behavior observed in the materials of the preceding section. Crack de￾Fig. 3. Push-out curves according to the roughness interface model, with parameters chosen to fit the experimental curve of Fig. 1, assuming that (i) initial cracking is governed by the roughness parameters h 4 50 nm, d 4 50 nm, and G 4 28 J/m2 and (ii) subsequent cracking is governed by h 4 50 nm, d 4 2.2 mm, and G 4 28 J/m2 ; the transition between the two conditions is indicated by the arrowed line. 1884 Journal of the American Ceramic Society—Kerans et al. Vol. 81, No. 7