l152 Journal of the American Ceramic Society--Sauder and Lamon Vol. 90. No. 4 1/T巛K-1 Stress(MPa) 52E045.6E0460E0464E0468E0472E04 1.00E-05 4 SA3(2)fiber(3DD Mpa) 10Eo6·s31eoco 00E-06 100E07 1.0E-08 D0E08 鲤 1364 kJ. mo 370 KJ. mor 1.0E09 16001500°1400°1300 200° Fig 9. Logarithmic plot of strain rate versus applied stress for SA3(2) Fig 8. Steady-state creep rate versus reciprocal temperature for SA3 The apparent activation energy determined on Hi-Nicalon S (770 kJ/mol, Table In is consistent with that corresponding to diffusion of carbon within grains(840 kJ/mor--) It seems too main difference between both SA3 fibers lies in the amount of low for silicon diffusion(910 kJ/mol) Diffusion within grains free carbon, this result indicates that secondary creep was not could be related to the feature of the microstructure. As men- influenced by the amount of carbon. tioned in section Ill (1), grain boundaries were not clearly de- Determination of stress exponent by fitting Eq(5)to creep fined. It may be thought that they cannot be preferred paths fo rates determined at different stresses(Fig 9) yields stress expo- diffusion. nents of 2.5 for all the fibers. It is worth mentioning that plots of The activation energy determined for SA3 fibers(370 kJ /mol train rates determined under various temperatures and applied Table ID)is close to that corresponding to diffusion of Al stresses(Figs.8-10), and creep parameters reported in Table ll grain boundaries in polycrystalline SiC (360-460 kJ/molss-y do not exhibit a noticeable discrepancy, which demonstrates the This suggests the contribution of diffusion of Al at grain bound pertinence of experimental work. ies. It is also close to half of that corresponding to diffusion of In polycrystalline SiC ceramics, n= I and creep results from or Si within grains for Hi-Nicalon S(Table ID). Hence, the diffusional processes either at grain boundaries(Coble creep. contribution of diffusion of C or Si at grain boundaries canno m=3)or within grains( Nabarro Herring creep m=2). A discarded activation energy corresponds to Nabarro Herring creep. The Figure 10 compares creep rates for Hi-Nicalon S and SA3(2) creep rate is very sensitive to grain size(Eq. 5): m=2 or 3. fibers under 850 MPa. It appears unambiguously that creep de Identical trends have been reported for both a- and B-sic formations are larger in SA3 fibers, despite the presence of larger Diffusional creep in polycrystalline SiC occurs by diffusion of rains. This trend can be attributed to the presence of Al. Alu minum is known to favor diffusion at a high temperature. This or by diffusion of impurities at grain boundaries. There is a causes an increase in the diffusion coefficient d(Eq. (5)). At very limited amount of data on diffusion of carbon and silicon temperatures above 1500oC, the SA3 fiber becomes more creep elements in Sic. There is no consensus about the diffusion rate. resistant than hi-Nicalon s Depending on the author, the diffusion rate of carbon could be 100 times faster than that of silicium, 26, 27 or it could be slower. 28 Activation energies for diffusion of carbon and silicium within (5) Tertiary Creep Stage igures 7 and ll clearly indicate that, at very high temperatures, Fi been estimated to be 84029 and 910 kJ/mol, 30 respectively. It the creep rate accelerates into a tertiary stage. In order to iden- is worth keeping in mind that creep of polycrystalline Sic tify the creep mechanism, tests were interrupted during the third cannot be controlled by dislocation moti <1700.C.31,32 For Nicalon SiC fibers tested in CO nel.the scanning electron microscopy. Figures 1l and 12 show the typ- activation energy was found to be consistent with the activation ical microstructure of fibers that was observed. It consists of two energies of thermally activated viscous flow of glasses at high distinct parts: the core and an annular region. The microstruc- ure of the core was unaffected, whereas Auger analysis showed between 2 and 3 that have been determined fo low oxygen content SiC fibers" correspond to grain-bound- 1/T(K-) ary sliding in the absence of a glassy phase. 52E-0456E046.0E-0464E-0 68E-0472E04 Creep of the fibers of this study (na 2.5)may be attributed 1.00E405 重理理理理 grain-boundary sliding, without grain elongation and glassy phase(Rachinger mechanism). In polycrystalline ceramics accommodation results from diffusion and fold formation at 1.00E-06 riple junctions. In SiC fibers, it probably involves carbon deformation 00E07 Table IL. Steady-State Creep Parameters 1.00E08 range(C) o(k/mor) 1.00E09 SA3(1 l150-1500 0-370)2.5(2.35-2.6) 1400°1300%1200° SA3(2) l150-1500 370)2.5(2.3-2.6 Hi-Nicalon s 1300-1500 770)2.6(2.42.9) state creep rate versus reciprocal temperature for SA3 and Hi- Nicalon S fibers under a stress of 850 MPamain difference between both SA3 fibers lies in the amount of free carbon, this result indicates that secondary creep was not influenced by the amount of carbon. Determination of stress exponent by fitting Eq. (5) to creep rates determined at different stresses (Fig. 9) yields stress expo￾nents of 2.5 for all the fibers. It is worth mentioning that plots of strain rates determined under various temperatures and applied stresses (Figs. 8–10), and creep parameters reported in Table II do not exhibit a noticeable discrepancy, which demonstrates the pertinence of experimental work. In polycrystalline SiC ceramics, n 5 1 and creep results from diffusional processes either at grain boundaries (Coble creep, m 5 3) or within grains (Nabarro Herring creep m 5 2). A high￾activation energy corresponds to Nabarro Herring creep. The creep rate is very sensitive to grain size (Eq. 5): m 5 2 or 3. Identical trends have been reported for both a- and b-SiC. Diffusional creep in polycrystalline SiC occurs by diffusion of carbon22 or silicon23 at grain boundaries22,23 or within grains,24 or by diffusion of impurities at grain boundaries.25 There is a very limited amount of data on diffusion of carbon and silicon elements in SiC. There is no consensus about the diffusion rate. Depending on the author, the diffusion rate of carbon could be 100 times faster than that of silicium,26,27 or it could be slower.28 Activation energies for diffusion of carbon and silicium within grains in b-SiC made via chemical vapor deposition have been estimated to be 84029 and 910 kJ/mol,30 respectively. It is worth keeping in mind that creep of polycrystalline SiC cannot be controlled by dislocation motions at temperatures o17001C.31,32 For Nicalon SiC fibers tested in CO, n1, the activation energy was found to be consistent with the activation energies of thermally activated viscous flow of glasses at high temperatures.33 n exponents between 2 and 3 that have been determined for low oxygen content SiC fibers34–36 correspond to grain-bound￾ary sliding in the absence of a glassy phase.37 Creep of the fibers of this study (n2.5) may be attributed to grain-boundary sliding, without grain elongation and glassy phase (Rachinger mechanism). In polycrystalline ceramics, accommodation results from diffusion and fold formation at triple junctions.37 In SiC fibers, it probably involves carbon deformation. The apparent activation energy determined on Hi-Nicalon S (770 kJ/mol, Table II) is consistent with that corresponding to diffusion of carbon within grains (840 kJ/mol28–29). It seems too low for silicon diffusion (910 kJ/mol30). Diffusion within grains could be related to the feature of the microstructure. As men￾tioned in section III (1), grain boundaries were not clearly de- fined. It may be thought that they cannot be preferred paths for diffusion. The activation energy determined for SA3 fibers (370 kJ/mol, Table II) is close to that corresponding to diffusion of Al at grain boundaries in polycrystalline SiC (360–460 kJ/mol38–40). This suggests the contribution of diffusion of Al at grain bound￾aries. It is also close to half of that corresponding to diffusion of C or Si within grains for Hi-Nicalon S (Table II). Hence, the contribution of diffusion of C or Si at grain boundaries cannot be discarded. Figure 10 compares creep rates for Hi-Nicalon S and SA3 (2) fibers under 850 MPa. It appears unambiguously that creep de￾formations are larger in SA3 fibers, despite the presence of larger grains. This trend can be attributed to the presence of Al. Alu￾minum is known to favor diffusion at a high temperature. This causes an increase in the diffusion coefficient D (Eq. (5)). At temperatures above 15001C, the SA3 fiber becomes more creep resistant than Hi-Nicalon S. (5) Tertiary Creep Stage Figures 7 and 11 clearly indicate that, at very high temperatures, the creep rate accelerates into a tertiary stage. In order to iden￾tify the creep mechanism, tests were interrupted during the third stage and the cross section of the fibers was examined using scanning electron microscopy. Figures 11 and 12 show the typ￾ical microstructure of fibers that was observed. It consists of two distinct parts: the core and an annular region. The microstruc￾ture of the core was unaffected, whereas Auger analysis showed 1600 °C 1500 °C 1400 °C 1300 °C 1200 °C 1.00E-09 1.00E-08 1.00E-07 1.00E-06 1.00E-05 5.2E-04 5.6E-04 6.0E-04 6.4E-04 6.8E-04 7.2E-04 1/T (K ) 370 kJ.mol 372 kJ.mol 364 kJ.mol ε (s ) . Fig. 8. Steady-state creep rate versus reciprocal temperature for SA3 (2) fiber. Table II. Steady-State Creep Parameters Fibers Temperature range (1C) Activation energy Q (kJ/mol
1 ) n SA3 (1) 1150–1500 370 (360–370) 2.5 (2.35–2.6) SA3 (2) 1150–1500 370 (360–370) 2.5 (2.3–2.6) Hi-Nicalon S 1300–1500 770 (750–770) 2.6 (2.4–2.9) 1.0E-09 1.0E-08 1.0E-07 1.0E-06 1.0E-05 100 Stress (MPa) 1000 ε (s ) . Fig. 9. Logarithmic plot of strain rate versus applied stress for SA3 (2) fiber. 1400 °C 1300 °C 1200 °C 1.00E-09 1.00E-08 1.00E-07 1.00E-06 1.00E-05 5.2E-04 5.6E-04 6.0E-04 6.4E-04 6.8E-04 7.2E-04 1/T (K ) ε (s ) . Fig. 10. Steady-state creep rate versus reciprocal temperature for SA3 (2) and Hi-Nicalon S fibers under a stress of 850 MPa. 1152 Journal of the American Ceramic Society—Sauder and Lamon Vol. 90, No. 4