J.Am. Ceram.Soe.84I]2451-53(2001) journa Experimental Observations of High Fracture Resistance in Silicon Carbide/Alumina Laminated Composites Ramanthan Krishnamurthy and Brian W. Sheldon Division of Engineering, Brown University, Providence, Rhode Island 02912 Model laminated composites were fabricated with porou experiments performed here. In this work, we use a porous-Al2O3 Al2O3 interfaces between SiC bars. The porous Al2O, was interface between SiC bars. This system inherently possesses deposited using an aerosol spray deposition technique, and the residual thermal stresses because of the coefficient of thermal andwich specimen was fabricated by hot pressing. Residual expansion(CTE)mismatch between Al, O, and SiC, which should thermal stresses were present in the interface because of the affect the interfacial fracture resistances measured and the com difference in the coefficients of thermal expansion of SiC and petition between crack deflection along the interface and crack Al,O3 Crack deflection was observed with measured interfa- penetration through the interface ial fracture resistances that were considerably higher than the Changes in the He-Hutchinson criterion because of in-plane deflection threshold predicted by the He-Hutchinson criterion and normal residual stresses have been studied by he et al. 7 This Examination of the fracture surface revealed a tortuous crack analysis also quantifies the effects of elastic mismatch using the path and significant crack-flaw interaction parameter a = DE Er)(i.e, Dundurs paramete representing elastic mismatch between the layers I and 2, where and E, represent the plane strain elastic moduli of materials I an L. Introduction 2, respectively ). The He et al. results show that enhanced fracture ROGREsS in technology, such as gas turbines in aircraft engines, uires the development of new and superior hi in layered materials with significant in-plane residual stresses, temperature materials. Ceramic materials with their high meltin interposing thin layers of a material with high a and low CTI points and excellent creep resistances are the obvious candidates (compared with the matrix material ). Based on this analysis, these enhancements are not expected to occur for the SiC/porous-ALO for high-temperature applications. However, ceramic materials system (low a and high CTE). The experimental results tha matrix composites have significantly higher fracture resistances we present in this communication contradict these theoretical compared with monolithic ceramics and seem to be ideal candi- predicti dates for use in high-temperature applications For improved fracture resistance in brittle-matrix composite IL. Experimental Procedure the use of a weak interface that promotes crack deflection is cessary. The He-Hutchinson criterion stipulates the The Sic bars(CVD SiC, rohm and Haas Advanced materials, requirement for crack deflection in terms of the ratio of the Woburn, MA)were machined to dimensions of 2 mm X 4 mm X interfacial fracture resistance Ii to the substrate fracture resistance 70 mm, and the surface of the bars was lapped with I um diameter Tf For a composite with no elastic mismatch between the layers, grit to a surface finish of Ra 0.05 um. The interfacial layer of porous Al,O3 was deposited on the Sic bars using an aerosol spray The earliest ceramic composites used bn or carbon as inter deposition method that is described elsewhere. Very fine(37nm s. However, these materials are prone to oxidation at higl sized particles in the green state)Al,O, powder(aluminum oxide temperatures. Porous-oxide layers seem to be an attractive alter- Nanophase Technologies Corp, Burr Ridge, IL) was used to create native and have been successfully demonstrated as effective the interfacial layer. A composite sandwich with the dimensions 4 interface layers in laminated ceramic composites. "4 The use of mm X 4 mm X 70 mm was formed by placing an identical SiC bar porous layers exploits the fact that as little as 10% porosity can on the deposited bar. a die machined to tolerances of +25 uI cause a considerable reduction in fracture toughness. Earlier work from high-quality graphite( Poco ADF-10Q, Poco Graphite, Inc with a laminated system consisting of porous-Al,O, interfaces Decatur, TX) was used to align the two bars of the sandwich. A between Al,O, bars showed markedly improved fracture resis mating graphite platen was placed atop the sandwich. The sand tance for these composites as compared with monolithic Al,O,.. wich was hot-pressed for 4 h under a vacuum of 10-6 torr(1.3X Also, the interfacial fracture resistances measured in the previous 10-4 Pa)at the various temperature and pressure combinations work are consistent with the he-Hutchinson criterion listed in Table I. The thickness of the interfacial layer after hot However, Al,, is known to creep significantly at temperatures pressing was -2 um Based on the elastic constants and CTEs of of -1200oC In comparison, SiC has much better creep resistance SiC and Al2O3, the biaxial residual thermal stress in the plane of and, thus, has been selected as the model matrix material for the the interphase was determined to be 2.5 GPa and tensile in nature (based on g= MA(CTE)AT, where M is the biaxial modulus, A(CTE) the CTE mismatch, and AT the difference between the hot-pressing and ambient temperatures) D. B. Marshalk--contributing editor Figure I is a schematic representation of the test geometry. A Vickers indentation technique was used to introduce precrack that were later propagated through the upper matrix bar under three- oint loading. The details of the precracking procedure are Manuscript No. 188129, Received November 27, 2000, approved June 26, 2001 described elsewhere. In the three-point bend test, the composite bar was unloaded as soon as the main crack reached the interface Member, American Ceramic Society and arrested. The specimen was then tes
Experimental Observations of High Fracture Resistance in Silicon Carbide/Alumina Laminated Composites Ramanthan Krishnamurthy and Brian W. Sheldon* Division of Engineering, Brown University, Providence, Rhode Island 02912 Model laminated composites were fabricated with porousAl2O3 interfaces between SiC bars. The porous Al2O3 was deposited using an aerosol spray deposition technique, and the sandwich specimen was fabricated by hot pressing. Residual thermal stresses were present in the interface because of the difference in the coefficients of thermal expansion of SiC and Al2O3. Crack deflection was observed with measured interfacial fracture resistances that were considerably higher than the deflection threshold predicted by the He–Hutchinson criterion. Examination of the fracture surface revealed a tortuous crack path and significant crack–flaw interaction. I. Introduction PROGRESS in technology, such as gas turbines in aircraft engines, requires the development of new and superior hightemperature materials. Ceramic materials with their high melting points and excellent creep resistances are the obvious candidates for high-temperature applications. However, ceramic materials have low fracture toughness that limit their potential. Ceramicmatrix composites have significantly higher fracture resistances compared with monolithic ceramics and seem to be ideal candidates for use in high-temperature applications. For improved fracture resistance in brittle-matrix composites, the use of a weak interface that promotes crack deflection is generally necessary.1 The He–Hutchinson criterion stipulates the requirement for crack deflection in terms of the ratio of the interfacial fracture resistance i to the substrate fracture resistance f . 2 For a composite with no elastic mismatch between the layers, the threshold value of this ratio is 0.25. The earliest ceramic composites used BN or carbon as interfaces. However, these materials are prone to oxidation at high temperatures. Porous-oxide layers seem to be an attractive alternative and have been successfully demonstrated as effective interface layers in laminated ceramic composites.3,4 The use of porous layers exploits the fact that as little as 10% porosity can cause a considerable reduction in fracture toughness.5 Earlier work with a laminated system consisting of porous-Al2O3 interfaces between Al2O3 bars showed markedly improved fracture resistance for these composites as compared with monolithic Al2O3. 6 Also, the interfacial fracture resistances measured in the previous work are consistent with the He–Hutchinson criterion. However, Al2O3 is known to creep significantly at temperatures of 1200°C. In comparison, SiC has much better creep resistance and, thus, has been selected as the model matrix material for the experiments performed here. In this work, we use a porous-Al2O3 interface between SiC bars. This system inherently possesses residual thermal stresses because of the coefficient of thermal expansion (CTE) mismatch between Al2O3 and SiC, which should affect the interfacial fracture resistances measured and the competition between crack deflection along the interface and crack penetration through the interface. Changes in the He–Hutchinson criterion because of in-plane and normal residual stresses have been studied by He et al.7 This analysis also quantifies the effects of elastic mismatch using the parameter (E 1 – E 2)/(E 1 E 2) (i.e., Dundurs parameter representing elastic mismatch between the layers 1 and 2, where E 1 and E 2 represent the plane strain elastic moduli of materials 1 and 2, respectively). The He et al. results show that enhanced fracture resistances compared with the monolithic material can be obtained in layered materials with significant in-plane residual stresses, by interposing thin layers of a material with high and low CTE (compared with the matrix material). Based on this analysis, these enhancements are not expected to occur for the SiC/porous-Al2O3 system (low and high CTE). The experimental results that we present in this communication contradict these theoretical predictions.7 II. Experimental Procedure The SiC bars (CVD SiC, Rohm and Haas Advanced Materials, Woburn, MA) were machined to dimensions of 2 mm 4 mm 70 mm, and the surface of the bars was lapped with 1 m diameter grit to a surface finish of Ra 0.05 m. The interfacial layer of porous Al2O3 was deposited on the SiC bars using an aerosol spray deposition method that is described elsewhere.8 Very fine (37 nm sized particles in the green state) Al2O3 powder (aluminum oxide, Nanophase Technologies Corp., Burr Ridge, IL) was used to create the interfacial layer. A composite sandwich with the dimensions 4 mm 4 mm 70 mm was formed by placing an identical SiC bar on the deposited bar. A die machined to tolerances of 25 m from high-quality graphite (Poco ADF-10Q, Poco Graphite, Inc., Decatur, TX) was used to align the two bars of the sandwich. A mating graphite platen was placed atop the sandwich. The sandwich was hot-pressed for 4 h under a vacuum of 10 6 torr (1.3 10–4 Pa) at the various temperature and pressure combinations listed in Table I. The thickness of the interfacial layer after hot pressing was 2 m. Based on the elastic constants and CTEs of SiC and Al2O3, the biaxial residual thermal stress in the plane of the interphase was determined to be 2.5 GPa and tensile in nature (based on M (CTE) T, where M is the biaxial modulus, (CTE) the CTE mismatch, and T the difference between the hot-pressing and ambient temperatures). Figure 1 is a schematic representation of the test geometry. A Vickers indentation technique was used to introduce precracks that were later propagated through the upper matrix bar under threepoint loading. The details of the precracking procedure are described elsewhere.8 In the three-point bend test, the composite bar was unloaded as soon as the main crack reached the interface and arrested. The specimen was then tested under four-point D. B. Marshall—contributing editor Manuscript No. 188129. Received November 27, 2000; approved June 26, 2001. Supported by the MRSEC Program of the National Science Foundation under Award No. DMR-0079964. *Member, American Ceramic Society. 2451 journal J. Am. Ceram. Soc., 84 [10] 2451–53 (2001)
2452 Communications of the American Ceramic Societ Vol 84. No 10 Table L. Selected T / T Values for Al, O3 Temperature material (J/m) fracture test 172A2O314.8 Deflected,025±00212 Deflected0.11±0.02 1400 10.4 SiC 15.4 Deflected 0.63 + 0.11 t-- alumina(1300 c) 1400 10.4 SiC 15.0 Deflected 0.62+ 0.11 1400 Broke through sic(1400c) eady-state load from boint bend test was then con- erted nterfacial resistance using a closed form Hotpressing pressure II. Results and discussion Fig. 2. Interfacial fracture resistance (in J/m-)as a function of hot- pressing pressure(in MPa) for various processing temperatures. Values of The measured fracture resistances are plotted as a function of the Al2O, system that has no elastic mismatch are plotted for comparison. hot-pressing pressure in Fig. 2, and the results of the fracture experiments are summarized in Table I. For comparison, the maximum fracture resistance value at which crack deflection wa interfacial fracture resistance of 15.0 Jm- was obtained for another observed in an Al,O, /porous-Al, O, interphase system is also specimen processed under the same conditions. This value is listed(obtained at hot-pressing temperature and pressure of within experimental error and confirms the validity of the result. A 1300.C and 17.2 MPa, respectively ) In this Al,,,O, high degree of reproducibility with this type of specimen also has system, crack penetration through the interface onto the lower bar been observed was observed when the interfacial fracture resistance was in- Several other issues need to be considered in applying the creased to a slightly higher value by processing at the same He-Hutchinson criterion to the experimental observations. The temperature and a higher pressure of 20.8 MPa. This set of basic He-Hutchinson derivation is based on energy release rates ocessing parameters was used as preliminary processing condi- for a homogeneous specimen with no separate layer to characterize tions for the work conducted here. As shown in Fig. 2, the the thin interphase. The effect of the elastic mismatch of a thin SiC/porous-Al,O, specimens exhibited lower interfacial fracture interlayer has been analyzed for the limiting case where the film resistances than the AlO,/porous-Al2O, specimens that were thickness asymptotically goes to zero. For the small elastic ocessed under identical conditions. The easier crack deflection mismatch of o=0.06 in our system, this effect leads to a slightly observed with SiC bars in our work(compared with the Al,O,/ smaller threshold for the energy release rate ratio, which rous-Al,O, system) indicated that the interfacial bonding be- contrast with the experimental observations. (This value of a was tween the SiC and porous Al,O3 was relatively weak. calculated using the elastic constants for dense Al,O3. With The monolithic CVD-SiC used as the matrix material has roughly 20% porosity for the interlayer in our experiments, o is fracture resistances of 20.6-29 7 J/m- with an average value o slightly larger if the elastic modulus varies linearly with porosity. 25.1 J/m. Taking one-fourth of the fracture resistance of the The effect of finite interphase thickness on the deflection threshold CVD-SiC as the threshold value for the interfacial fracture also has been explored by treating the laminate as a three-layer criterion for crack deflection in the absence of residual stress creased to higher r. this effect flattens out after some decrease in would be between 5.2 and 7.4 J/m- with an average value of 6.3 the interphase thickness, and, in addition, is significant only for a J/m. For the actual system, where there are in-plane tensile high negative value of a(as opposed to the positive value of a in residual thermal stresses in the porous-Al,O, layer, the threshold our syster values for Ii would be slightly lower than in the unstressed system, Another issue is that the theoretical predictions were formulated according to predictions made by He et al. In contrast, the a pupative crack at the interface, while the measured fracture measured fracture resistance of 15.4 J/m- obtained at a tempera- resistance values were obtained for deflected cracks that can be ture/pressure combination of 1400.C and 10.4 MPa was signifi- several millimeters long. The He-Hutchinson criterion is in cantly higher than the predicted threshold values. A similar excellent agreement with the analogous data obtained with Al O porous-Al2O3 laminates, which suggests that there are no signif cant differences between the initial crack and the longer, deflected crack. These previous specimens were fabricated and tested in the same way as the Sic/porous-Al2O3 specimens, but this does not guarantee that the measured fracture resistance values accurately represent the behavior of a pupative crack in this case as well especially because some differences in the behavior of the inter- ohase were observed in these two systems. R-curve effects were not observed with either of these systems (i.e, th 0.002mm constant during the steady-state crack growth that used to obtain the values in Table I and Ref. 8). Thus fference between a pupative crack and the measured results would hav Fig. 1. Experimental setup to interfacial fracture res ape detection with this measurement under four-point bending. Main crack was initiated from Vickers strikes on The basic mechanics analysis for a homogeneous interface lay the lower bar and was subsequently arrested in the interface. predicts that crack deflection occurs at the upper interface of our
bending. This procedure is also described elsewhere.8,9 The steady-state load from the four-point bend test was then converted to an interfacial fracture resistance using a closed form solution.9,10 III. Results and Discussion The measured fracture resistances are plotted as a function of hot-pressing pressure in Fig. 2, and the results of the fracture experiments are summarized in Table I. For comparison, the maximum fracture resistance value at which crack deflection was observed in an Al2O3/porous-Al2O3 interphase system is also listed (obtained at hot-pressing temperature and pressure of 1300°C and 17.2 MPa,8 respectively). In this Al2O3/porous-Al2O3 system, crack penetration through the interface onto the lower bar was observed when the interfacial fracture resistance was increased to a slightly higher value by processing at the same temperature and a higher pressure of 20.8 MPa. This set of processing parameters was used as preliminary processing conditions for the work conducted here. As shown in Fig. 2, the SiC/porous-Al2O3 specimens exhibited lower interfacial fracture resistances than the Al2O3/porous-Al2O3 specimens that were processed under identical conditions. The easier crack deflection observed with SiC bars in our work (compared with the Al2O3/ porous-Al2O3 system) indicated that the interfacial bonding between the SiC and porous Al2O3 was relatively weak. The monolithic CVD-SiC used as the matrix material has fracture resistances of 20.6–29.7 J/m2 with an average value of 25.1 J/m2 . 11 Taking one-fourth of the fracture resistance of the CVD-SiC as the threshold value for the interfacial fracture resistance, the maximum values allowed by the He–Hutchinson criterion for crack deflection in the absence of residual stress would be between 5.2 and 7.4 J/m2 with an average value of 6.3 J/m2 . For the actual system, where there are in-plane tensile residual thermal stresses in the porous-Al2O3 layer, the threshold values for i would be slightly lower than in the unstressed system, according to predictions made by He et al. 7 In contrast, the measured fracture resistance of 15.4 J/m2 obtained at a temperature/pressure combination of 1400°C and 10.4 MPa was significantly higher than the predicted threshold values. A similar interfacial fracture resistance of 15.0 J/m2 was obtained for another specimen processed under the same conditions. This value is within experimental error and confirms the validity of the result. A high degree of reproducibility with this type of specimen also has been observed.6,8 Several other issues need to be considered in applying the He–Hutchinson criterion to the experimental observations. The basic He–Hutchinson derivation is based on energy release rates for a homogeneous specimen with no separate layer to characterize the thin interphase. The effect of the elastic mismatch of a thin interlayer has been analyzed for the limiting case where the film thickness asymptotically goes to zero.12 For the small elastic mismatch of 0.06 in our system, this effect leads to a slightly smaller threshold for the energy release rate ratio, which is in contrast with the experimental observations. (This value of was calculated using the elastic constants for dense Al2O3. With roughly 20% porosity for the interlayer in our experiments, is slightly larger if the elastic modulus varies linearly with porosity.) The effect of finite interphase thickness on the deflection threshold also has been explored by treating the laminate as a three-layer system.13 With thinner interphases, the deflection limit is increased to higher i . This effect flattens out after some decrease in the interphase thickness, and, in addition, is significant only for a high negative value of (as opposed to the positive value of in our system). Another issue is that the theoretical predictions were formulated for a pupative crack at the interface, while the measured fracture resistance values were obtained for deflected cracks that can be several millimeters long. The He–Hutchinson criterion is in excellent agreement with the analogous data obtained with Al2O3/ porous-Al2O3 laminates, which suggests that there are no significant differences between the initial crack and the longer, deflected crack.8 These previous specimens were fabricated and tested in the same way as the SiC/porous-Al2O3 specimens, but this does not guarantee that the measured fracture resistance values accurately represent the behavior of a pupative crack in this case as well, especially because some differences in the behavior of the interphase were observed in these two systems. R-curve effects were not observed with either of these systems (i.e., the load was constant during the steady-state crack growth that was used to obtain the values in Table I and Ref. 8). Thus, any difference between a pupative crack and the measured results would have to escape detection with this measurement. The basic mechanics analysis for a homogeneous interface layer predicts that crack deflection occurs at the upper interface of our Fig. 1. Experimental setup to measure interfacial fracture resistance under four-point bending. Main crack was initiated from Vickers strikes on the lower bar and was subsequently arrested in the interface. Fig. 2. Interfacial fracture resistance (in J/m2 ) as a function of hotpressing pressure (in MPa) for various processing temperatures. Values of the Al2O3 system that has no elastic mismatch are plotted for comparison.8 Table I. Selected i /f Values for Al2O3 † and SiC Composites Temperature (°C) Pressure (MPa) Substrate material Interfacial fracture resistance (J/m2 ) Result of fracture test i /f 1300 17.2 Al2O3 14.8 Deflected 0.25 0.02 1300 20.8 Al2O3 Broke through 1300 17.2 SiC 0.5 Deflected 1300 20.8 SiC 2.9 Deflected 0.11 0.02 1400 7.8 SiC 8.9 Deflected 0.36 0.06 1400 10.4 SiC 15.4 Deflected 0.63 0.11 1400 10.4 SiC 15.0 Deflected 0.62 0.11 1400 17.2 SiC Broke through † Reference 8. 2452 Communications of the American Ceramic Society Vol. 84, No. 10
October 2001 Communications of the American Ceramic Sociery 453 testing geometry (i.e, the second SIC/Al,O, interface that the that the threshold energy release rate ratio for cont deflection crack encounters). A micrograph of the fractured surface of the s-0.6.This threshold criterion can be applical pper Sic bar is shown in Fig. 3, where there is a substantial interphase, where a crack is likely to meet an acial flaw amount of Al,O, adhering to the relatively smooth SiC surface. However, the interfacial flaws in our system may not be large This suggests significant microstructural inhomogeneity in either enough to be treated as interfacial cracks. The effect of porosity on the porous interphase or at the SiC/Al,O3 interface. The tortuous ontinued crack deflection also has been studied in a similar fracture path indicates that the deflected crack is not moving manner, by treating the porosity as flaws smoothly along the upper interface. This effect should increase the fracture resistance of a deflected crack, which is potentially consistent with the experimentally measured fracture resistance IV. Conclusions values. Unfortunately, the effects of such microstructure distribu- tions on the deflection criterion are not clearly understood. A Model laminates were fabricated using thin, porous-AL2O3 already described, the Al2O/Al2O3 system that was investigated interphases between SiC substrates. The interfacial fracture resis previously has a porous microstructure that is similar to the tance in these systems was lower than that observed in interphase investigated here. However, the type of fracture surface AL,O,porous-AL2O3 system processed under identical conditions in Fig 3 was not observed. Thus, flaw population effects leading Experiments also showed that crack deflection occurred in speci mens where the interfacial fracture resistance significantly ex to the higher than expected T, values are more likely to be related ceeded the deflection limit defined by the He-Hutchinson crite to the SiC/porous- A2 O, interface, rather than to the bulk porous. rion. Porosity and flaw distributions along the interphase may have ignificant effects on the deflection criterion and a more detailed somewhat different because of the higher hot-pressing temperature examination of those issues may have important implications for and the possible presence of Sio,(native oxide on SiC) Work by Mammoli et al. suggests that the presence of flaws the design of brittle-matrix composites has a considerable effect on the crack deflection criterion. Based on their analysis, flaws increase the allowable energy release ratio, Acknowledgmen such that higher interfacial resistances allow crack deflection. This ay partly explain how interfacial fracture resistances above the We are grateful to Dr, Jit Goela of Rohm and Haas Advanced Materials for He-Hutchinson limit may lead to crack deflection. However, the providing the SiC substrate material, and to Professors Janet Rankin and David Green for their comments on the manuscript. effects of flaws are significant only when the crack tip gets close to the flaw, and energy release rate ratios well in excess of the He-Hutchinson threshold can lead to crack deflection only when References the crack meets a flaw at the interface. Residual stress effects ior of Ceramic Matrix such crack-flaw interactions and their implications with regard to Composites Metall. mater.,37012567-83(1989) the competition between crack deflection and penetration are Dissimilar Elastic Materials, "Int J. Solids Struct, 25 19)1053-67(1989) problems that have not been addressed For the specific case of the 3M. J. O'Brien, F M. Capaldi, crack meeting a flaw at the interface, the flaw can be considered as an extension of the crack onto the interface With a homogeneous K.S. Blanks, A Kristoffersson, E Carlstrom, and w.J. Clegg, "Crack Deflectio interface, where a defect lies along the interface, it has been shown in Ceramic Laminates Using Porous Interlayers,J. Ear. Ceram. Soc., I8 L.A. Simpson, "Effect of Microstructure on Measurements of Fracture Energy of Al2O3,J Am Ceram Soc., 56 117-11(19 gh-T. O'Brien, "Fabrication of a Tailored Oxidation-Resistant Interface and W. Hutchinson, "Crack Deflection at an Interface between Dissimilar Elastic Materials: Role of Residual Stresses, Int J Solids Struct. 31口243443-55(199 SM. J. O'Brien and B. W. Sheldon, "Porous Alumina Co Fracture Resistance for Alumina Composites, ".Am. Ceram. Soc., 82[12]3567- PP. G Charalambides, J. Lund, A. G. Evans, and R. M. McMeeking,"A Test P G. Charalambides, H. C Cao, J. Lund, and A G. Evans, " Development of a Test Method for Measuring the Mixed Mode Fracture Resistance,J. App/ Mech, 8 269-83(1989 IM. A. Pickering, R. L. Taylor, J. T. Keeley, and G.A. Graves,"Chemically Vapor Deposited Silicon Carbide (SiC)for Optical Applications, Nucl. Instrum. Methods,A291,95-110(1990) 协 er.Si.Eg,A07,135-43(1989) Y F. Liu, Y. Tanaka, and C. Masuda, "Debonding Mechanisms in the Presence 078“15001Nmc0招 s," dcta metall. Mater,465]5237-47(1998 mmoli, A. L. Graham, I. E. Reimanis, and D. L. Tullock, "The Effect of Flaws on the Propagation of Cracks at Bimaterial Interfaces, Acta Metall. Mater Fig 3. Micrograph of the fracture surface on the upper SiC interface M. Y. He and J. w. Hutchinson."Kinking of a Crack Out of an Interface contact with the porous interphase J. AppL. Mech.,56,270-78(1989)
testing geometry (i.e., the second SiC/Al2O3 interface that the crack encounters). A micrograph of the fractured surface of the upper SiC bar is shown in Fig. 3, where there is a substantial amount of Al2O3 adhering to the relatively smooth SiC surface. This suggests significant microstructural inhomogeneity in either the porous interphase or at the SiC/Al2O3 interface. The tortuous fracture path indicates that the deflected crack is not moving smoothly along the upper interface. This effect should increase the fracture resistance of a deflected crack, which is potentially consistent with the experimentally measured fracture resistance values. Unfortunately, the effects of such microstructure distributions on the deflection criterion are not clearly understood. As already described, the Al2O3/Al2O3 system that was investigated previously has a porous microstructure that is similar to the interphase investigated here. However, the type of fracture surface in Fig. 3 was not observed.8 Thus, flaw population effects leading to the higher than expected i values are more likely to be related to the SiC/porous-Al2O3 interface, rather than to the bulk porousAl2O3 structure. It is also possible that the bulk porous structure is somewhat different because of the higher hot-pressing temperature and the possible presence of SiO2 (native oxide on SiC). Work by Mammoli et al.14 suggests that the presence of flaws has a considerable effect on the crack deflection criterion. Based on their analysis, flaws increase the allowable energy release ratio, such that higher interfacial resistances allow crack deflection. This may partly explain how interfacial fracture resistances above the He–Hutchinson limit may lead to crack deflection. However, the effects of flaws are significant only when the crack tip gets close to the flaw, and energy release rate ratios well in excess of the He–Hutchinson threshold can lead to crack deflection only when the crack meets a flaw at the interface. Residual stress effects on such crack–flaw interactions and their implications with regard to the competition between crack deflection and penetration are problems that have not been addressed. For the specific case of the crack meeting a flaw at the interface, the flaw can be considered as an extension of the crack onto the interface. With a homogeneous interface, where a defect lies along the interface, it has been shown that the threshold energy release rate ratio for continued deflection is 0.6.15 This threshold criterion can be applicable with a porous interphase, where a crack is likely to meet an interfacial flaw. However, the interfacial flaws in our system may not be large enough to be treated as interfacial cracks. The effect of porosity on continued crack deflection also has been studied in a similar manner, by treating the porosity as flaws.4 IV. Conclusions Model laminates were fabricated using thin, porous-Al2O3 interphases between SiC substrates. The interfacial fracture resistance in these systems was lower than that observed in an Al2O3/porous-Al2O3 system processed under identical conditions. Experiments also showed that crack deflection occurred in specimens where the interfacial fracture resistance significantly exceeded the deflection limit defined by the He–Hutchinson criterion. Porosity and flaw distributions along the interphase may have significant effects on the deflection criterion, and a more detailed examination of those issues may have important implications for the design of brittle-matrix composites. Acknowledgment We are grateful to Dr. Jit Goela of Rohm and Haas Advanced Materials for providing the SiC substrate material, and to Professors Janet Rankin and David Green for their comments on the manuscript. References 1 A. G. Evans and D. B. Marshall, “The Mechanical Behavior of Ceramic Matrix Composites,” Acta Metall. Mater., 37 [10] 2567–83 (1989). 2 M. Y. He and J. W. Hutchinson, “Crack Deflection at an Interface between Dissimilar Elastic Materials,” Int. J. Solids Struct., 25 [9] 1053–67 (1989). 3 M. J. O’Brien, F. M. Capaldi, and B. W. Sheldon, “A Layered Alumina Composite Tested at High Temperature in Air,” J. Am. Ceram. Soc., 83 [12] 3033–40 (2000). 4 K. S. Blanks, A. Kristoffersson, E. Carlstrom, and W. J. Cleggs, “Crack Deflection in Ceramic Laminates Using Porous Interlayers,” J. Eur. Ceram. Soc., 18 [13] 1945–51 (1998). 5 L. A. Simpson, “Effect of Microstructure on Measurements of Fracture Energy of Al2O3,” J. Am. Ceram. Soc., 56 [1] 7–11 (1973). 6 M. J. O’Brien, “Fabrication of a Tailored Oxidation-Resistant Interface and High-Temperature Testing of a Laminated Composite”; Ph..D. Thesis, Brown University, Providence, RI, 1998. 7 M. Y. He, A. G. Evans, and J. W. Hutchinson, “Crack Deflection at an Interface between Dissimilar Elastic Materials: Role of Residual Stresses,” Int. J. Solids Struct., 31 [24] 3443–55 (1994). 8 M. J. O’Brien and B. W. Sheldon, “Porous Alumina Coating with Tailored Fracture Resistance for Alumina Composites,” J. Am. Ceram. Soc., 82 [12] 3567–74 (1999). 9 P. G. Charalambides, J. Lund, A. G. Evans, and R. M. McMeeking, “A Test Specimen for Determining the Fracture Resistance of Bimaterial Interfaces,” J. Appl. Mech., 56, 77–82 (1989). 10P. G. Charalambides, H. C. Cao, J. Lund, and A. G. Evans, “Development of a Test Method for Measuring the Mixed Mode Fracture Resistance,” J. Appl. Mech., 8, 269–83 (1989). 11M. A. Pickering, R. L. Taylor, J. T. Keeley, and G. A. Graves, “Chemically Vapor Deposited Silicon Carbide (SiC) for Optical Applications,” Nucl. Instrum. Methods, A291, 95–110 (1990). 12Z. Suo and J. W. Hutchinson, “Sandwich Test Specimens for Measuring Interface Toughness,” Mater. Sci. Eng., A107, 135–43 (1989). 13Y. F. Liu, Y. Tanaka, and C. Masuda, “Debonding Mechanisms in the Presence of an Interphase in Composites,” Acta Metall. Mater., 46 [15] 5237–47 (1998). 14A. A. Mammoli, A. L. Graham, I. E. Reimanis, and D. L. Tullock, “The Effect of Flaws on the Propagation of Cracks at Bimaterial Interfaces,” Acta Metall. Mater., 43 [3] 1149–56 (1995). 15M. Y. He and J. W. Hutchinson, “Kinking of a Crack Out of an Interface,” J. Appl. Mech., 56, 270–78 (1989). Fig. 3. Micrograph of the fracture surface on the upper SiC interface that was in contact with the porous interphase. October 2001 Communications of the American Ceramic Society 2453