《复合材料 Composites》课程教学资源（学习资料）第五章 陶瓷基复合材料_j.1151-2916.1997.tb03056.x[1]

Joumal J Am Ceram Soc, 80[7]1812-20(1997) High-Temperature Transverse fracture Toughness of Nicalon-Fiber-Reinforced CAS-Il Glass-Ceramic Matrix Composite Ramazan Kahama Department of Chemical Engineering, King Fahd University of Petroleum and minerals Dhahran 31261, Saudi arabia John F. Mandell and Max C Deibert Department of Chemical Engineering, Montana State University, Bozeman, Montana Cracking parallel to the fibers in off-axis plies is usually the reinforced glass-ceramics, producing the carbon-rich inter- ng process has been associated with the(transverse)frac- bond strong enough for load transfer, yet weafer-matrix) ture toughness, defined by the critical strain energy release debond readily and allow fiber bridging during crack propaga ate, Gye. The measurement of Ge provides basic informa- tion(perpendicular to the fibers )resulting in high longitudinal tion about the transverse crack resistance. In this study, the fracture toughness in the composite. tility of the double torsion Dt) test technique to deter However, reductions of strength and strain to failure (in nine Gi in a glass-ceramic matrix composite(Nicalon/ the fiber direction) have been observed when testing at tem- CAS-Im at temperatures up to 1000.C has been demor peratures as low as 400oC. 10 In contrast to room-tempera temperature (as does the bulk matrix); however, no evi. pullout was observed at high temperatures. It has been pro- dence of an interphase oxidizing effect on crack growth posed that the transition from tough(notch- insensitive)behav (parallel to the fibers)could be found. The inevitable mis. ior at room temperature to brittle(notch-sensitive) behavior at significant matrix crack interactions with the fibers re- tion effects at interfaces exposed to the environment. 61 5 bridging the crack in the DT specimens, in contrast to the and/or increased fiber-matrix bond strength caused by ported for other geometries such as double cantilever beam The oxidation reaction is not initiated until the matrix cracks and flexure specimens. pon stressing(in the fiber direction), allowing penetration of high-temperature air. 12.13 The air then infiltrates the composite and attacks the low-strength carbon-rich fiber-matrix inter brittlement of the lot ERAMIc and glass-ceramic composites offer a great range of gitudinal fracture process..o t was reported that upon the oxi- utility on a high-temperature basis and based on other en- formed between the fiber and the matrix 14-l vironmental considerations such as oxidation resistance,ero- sion, and chemical attack from acids, etc. 2 The proper Inspite of the well-demonstrated embrittlement effect of ties of these composites depend on the combination of the reinforced glass-ceramic composites, there is not much study properties of starting materials and the fabrication proce- on the effects of temperature on fracture parallel to the fibers dure. By choosing high-strength, relatively high-modulus, continuous(or relatively long) fibers and incorporating them (transverse fracture). It was the objective of this study to in into matrices, high-strength and tough composites can be cre- vestigate the resistance of Nicalon-fiber-reinforced calcium ated if the fiber-matrix interphase is weak enough to divert aluminosilicate glass-ceramic matrix composite(Nicalon/CAS and/or bridge cracks. The ultimate strength of such a com- posite is then controlled by the in situ fiber strength, and a study of the transverse fracture toughness, the effects of high temperature air environment on fracture toughness, and also composite tensile strength well beyond that associated with the stress corrosion crack growth(if present) matrix cracking can be achieved. 3 The Nicalon-fiber reinforced glass-ceramic composites have met these require. ics is important since there are few engineering applications interphase during processing, and with no deterioration in fiber where transverse stresses are not encountered, 17, 18 The resis- tance to crack growth parallel to the fibers is especially a de- Nicalon has a unique nonstoichiometric chemistryis that ing parallel to the fibers is usually the initial form of damage in glass-ceramic matrix composites. Chemical reactions can occur multidirectional composites. 9 This cracking process has beer associated with: the(transverse) fracture toughness, defined by of energy per unit plate thickness per unit crack extension. Unstable(fast) crack growth should occur when the strain en ergy release rate, Gr, reaches the critical value, GIc, for the G. Evans--contributing editor material. Measurements of the strain energy release rate for fractures verify that the critical strain energy release rate is indeed a material parameter when fracture occurs exclusively by opening, Mode L, and when plane-strain conditions prevail 995: approved July 10, 1996. at the crack tip. 22,23 It is denoted by Gre, referring to failure in Mode I and under plain-strain conditions. Because of their

;;j ,;; 2 r.,. . .; ,.i . .. .._,...,. > .,.. J. Am. Cerarn. Soc., 80 I71 1812-20 (1597) High-Temperature Transverse Fracture Toughness of Nicalon-Fiber-Reinforced CAS-I I Glass-Ceramic Matrix Composite Ramazan Kahramant Department of Chemical Engineering, King Fahd University of Petroleum and Minerals, Dhahran 31261, Saudi Arabia John F. Mandell and Max C. Deibert Department of Chemical Engineering, Montana State University, Bozeman, Montana 59717 Cracking parallel to the fibers in off-axis plies is usually the initial form of damage in composite laminates. This crack￾ing process has been associated with the (transverse) frac￾ture toughness, defined by the critical strain energy release rate, GI,. The measurement of GI, provides basic inf'orma￾tion about the transverse crack resistance. In this study, the utility of the double torsion (DT) test technique to deter￾mine GI, in a glass-ceramic matrix composite (Nicalod CAS-11) at temperatures up to 1OOO"C has been demon￾strated. GI, did decrease moderately with increasing temperature (as does the bulk matrix); however, no evi￾dence of an interphase oxidizing effect on crack growth (parallel to the fibers) could be found. The inevitable mis￾alignment of fibers in the material was not very efficient at bridging the crack in the DT specimens, in contrast to the significant matrix crack interactions with the fibers re￾ported for other geometries such as double cantilever beam and flexure specimens. I. Introduction ERAMIC and glass-ceramic composites offer a great range of C utility on a high-temperature basis and based on other en￾vironmental considerations such as oxidation resistance, ero￾sion, and chemical attack from acids, etc.lV2 The proper￾ties of these composites depend on the combination of the properties of starting materials and the fabrication proce￾dure.' By choosing high-strength, relatively high-modulus, continuous (or relatively long) fibers and incorporating them into matrices, high-strength and tough composites can be cre￾ated if the fiber-matrix interphase is weak enough to divert and/or bridge cracks. The ultimate strength of such a com￾posite is then controlled by the in situ fiber strength, and a composite tensile strength well beyond that associated with matrix cracking can be a~hieved.~ The Nicalon-fiber￾reinforced glass-ceramic composites have met these require￾ments through the development of a carbon-rich fiber-matrix interphase during processing, and with no deterioration in fiber strength. Nicalon has a unique nonstoichiometric chemisw5 that makes it particularly suited to the development of high-strength glass-ceramic matrix composites. Chemical reactions can occur between the fiber and matrix during processing of Nicalon￾A. G. Evans-contributing editor Manuscript No. 192457. Received July 17. 1995; approved July 10. 1996. 'Author to whom correspondence should be addressed. reinforced glass-ceramics, producing the carbon-rich inter￾phase regi~n.~."~ The carbon-rich layer forms a (fiber-matrix) bond strong enough for load transfer, yet weak enough to debond readily and allow fiber bridging during crack propaga￾tion (perpendicular to the fibers) resulting in high longitudinal fracture toughness in the c0mposite.6*~ However, reductions of strength and strain to failure (in the fiber direction) have been observed when testing at tem￾peratures as low as 400°C.L0 In contrast to room-tempera￾ture fractures, a nearIy planar fracture surface with little fiber pullout was observed at high temperatures. It has been pro￾posed that the transition from tough (notch-insensitive) behav￾ior at room temperature to brittle (notch-sensitive) behavior at elevated temperatures is likely due to fiber strength degradation and/or increased fiber-matrix bond strength caused by oxida￾tion effects at interfaces exposed to the The oxidation reaction is not initiated until the matrix cracks upon stressing (in the fiber direction), allowing penetration of high-temperature air.12J3 The air then infiltrates the composite and attacks the low-strength carbon-rich fiber-matrix inter￾phase in such a way as to cause an embrittlement of the lon￾gitudinal fracture process.'S6 It was reported that upon the oxi￾dation of the interphase carbon layer a stronger silica bond is formed between the fiber and the matix.'"I6 Inspite of the well-demonstrated embrittlement effect of temperature on the longitudinal behavior of the Nicalon￾reinforced glass-ceramic composites, there is not much study on the effects of temperature on fracture parallel to the fibers (transverse fracture). It was the objective of this study to in￾vestigate the resistance of Nicalon-fiber-reinforced calcium alurninosilicate glass-ceramic matrix composite (NicalodCAS￾II) to crack growth parallel to the fibers. This included the study of the transverse fracture toughness, the effects of high￾temperature air environment on fracture toughness, and also the stress corrosion crack growth (if present). Studying the transverse properties of fiber-reinforced ceram￾ics is important since there are few engineering applications where transverse stresses are not en~ountered.'~.'~ The resis￾tance to crack growth parallel to the fibers is especially a de￾termining factor for the useful design stress range since crack￾ing parallel to the fibers is usually the initial form of damage in multidirectional composites.'9 This cracking process has been associated with the (transverse) fracture toughness, defined by the critical strain energy release rate, GIc.20*21 G,, has the units of energy per unit plate thickness per unit crack extension. Unstable (fast) crack growth should occur when the strain en￾ergy release rate, GI, reaches the critical value, G,,. for the material. Measurements of the strain energy release rate for fractures verify that the critical strain energy release rate is indeed a material parameter when fracture occurs exclusively by opening, Mode I, and when plane-strain conditions prevail at the crack tip.22,23 It is denoted by Gn, refemng to failure in Mode I and under plain-strain conditions. Because of their ' 1812

July 1997 High-Temperature Transverse fracture Toughness l813 limited plasticity, materials like ceramics and glasses The theoretical analysis of the dt technique assu atisfy the plar ditions as long as the thickness of the crack front is straight and orthogonal to the plane of the tative of the material 23 ough to be microstructurally represen the specimen is specimen. However, in most materials the crack front profile is significantly curved, with the furthest advance along the bot tom face of the specimen. 4 Because of this, Evans2suggests I. Test Technique rather than analytically. It is important that a constant-K(stress intensity factor The consensus among most researchers, 32 has been specimen be used so that continuous crack length measure crack velocity and crack length of interest is that of the ments are not required for data reduction, since high dge of the crack, i.e the furthest advanced bottom temperature tests will take place within a furnace. the doubl location transverse fracture toughness in this study, satisfies the con- specimen dimensions are approximately as follows: wit um Tait Fry and Garrett 3 have suggested that the opti stant-K requirement. The utility of the dt technique has been demonstrated in both ceramic and in composites. 20, 2It thicknesses complicate the analysis because of the significant has been shown that it gives results similar to those of other test contact stresses generated by the interaction of the torsion arms The original DT technique was suggested by Outwater and Another area of contention for the dt test configuration is Gerry and developed by Kies and Clark. 27 As illustrated in Fig. the length of crack over which the stress intensity may be I, the test specimen is a rectangular plate with a narrow notch considered to be independent of crack length(constant-K re or crack started in one end. The specimen is loaded in four- point bending across the notch, i. e, in"double torsion. " Thus lengths greater than 0. 55w, and ligaments(uncracked portion) the crack propagates axially along the specimen The crack is greater than 0. 65W. Tait et al found a consensus over severa researchers which indicated that at least the central one-third t one-half of the specimen (for L/w =3) should offer shear deformation for the crack to develop modes Il and III constant-K conditions Shetty and Virkar4 stated that the valid failure vanishes because of the symmetric double torsion"of region decreases as the en length-to-width ratio de- 山 pecimen with respect to the crac2 creases First, the pecimen complia ance, c (inverse of the against displacement, d, curve), is linearly related to the spec Il. Experimental Procedure men crack length, a. Second, the critical load, Pc, at which crack propagation is initiated, is independent of crack length, a The materials used in this study were unidirectional ([0]16) (constant-K specimen). Therefore, the fracture toughness, GIe Nicalon fiber/CAS-Il matrix composites. They were supplied ed from the compliance by Corning, Inc in the form of square plates, approximately I UsIn cm by 15 cm and about 3 mm thick. dt tests were conducted on an instron model 8562 servo- P? dc electric universal testing machine. Tests were performed with a Gle=2 da (1) 500 lb Le Bow load cell mounted below the standard 20 000 lb Instron load cell. This lower capacity load cell was used be where t is the specimen thickness. Both Pc and dida can be cause of its better resolution in the load range of interest. An accurately determined from the test results, as will be discussed Inconel rod was attached to the load cell for load application in the Experimental Section A flat compression table was mounted to the bottom grip to The compliance relationship with respect to the crack length, hold the DT fixture base. a photograph of the loading section dClda, can also be expressed analytically by 28 of the dt test system is shown in Fig. 2 The test fixture used for the dt test was similar to the one designed by Professor Leon Chuck from the University of Day ton. A close-up of the fixture, with a composite specime mounted, appears in Fig. 3. The fore-and-aft, and side-to-side modifying the lengt由 the two pivot rollers:0m如b w(r=1-0.6302(2t/W)+1.20(2 /W)e m/ww (3) the top portion plays no role during loading. The fixture, the base plate, and the loading rod were all made from Inconel, a nickel-based superalloy, for use at ter mperatures up to 1000C. The testing machine was equipped with a slot type, 2 Zone Model"D", fumace, mounted on rails, supplied by Instron Corp This furmace which was originally designed for tensile sts was enlarged axially by plates to create a box-shaped furnace positioned over the slot furnace. Internal dimensions of the box were 15 cm x 15 cm For those tests performed at high temperatures using the en larged furnace, the temperature was read by inserting a ther mocouple wire into the furnace next to the dt specimen DT test specimen blanks of 2. 18 cm x 6.70 cm were cut from the composite plates using a diamond saw and a diamond coated grinding wheel. Initial notches were cut into one side of the specimens using the 1 mm thick diamond saw. To deter mine the compliance vs crack length calibration, dc/da, six to mens were prepared and a different notch length was machined into each specimen. The introduced notch length Fig. 1. Double torsion test configuration ranged from 1. 80 to 5.08 cm

July 1997 High-Temperature Transverse Fracture Toughness 1813 limited plasticity, brittle materials like ceramics and glasses satisfy the plane-strain conditions as long as the thickness of the specimen is large enough to be microstructurally represen￾tative of the material.23 11. Test Technique It is important that a constant-K (stress intensity factor) specimen be used so that continuous crack length measure￾ments are not required for data reduction, since high￾temperature tests will take place within a furnace. The double torsion (DT) test method, which was utilized to measure the transverse fracture toughness in this study, satisfies the con￾stant-K requirement. The utility of the DT technique has been demonstrated in both ceramics2626 and in composites.20.21 It has been shown that it gives results similar to those of other test technique^.^^.^^ The original DT technique was suggested by Outwater and Gerry and developed by Kies and Clark.27 As illustrated in Fig. 1, the test specimen is a rectangular plate with a narrow notch or crack started in one end. The specimen is loaded in four￾point bending across the notch, i.e., in “double torsion.” Thus the crack propagates axially along the specimen. The crack is more open at the bottom face and less open at the top face of the specimen, defining Mode I fra~t~re.~~.~~.~~*~~ The necessary shear deformation for the crack to develop Modes 11 and 111 failure vanishes because of the symmetric “double torsion” of the specimen with respect to the crack plane.2’.28 The DT test has two major features.20.21,2627,29.30 First, the specimen compliance, C (inverse of the initial slope of load, P, against displacement, d, curve), is linearly related to the speci￾men crack length, a. Second, the critical load, P,, at which crack propagation is initiated, is independent of crack length, a (constant-K specimen). Therefore, the fracture toughness, GI,, can be computed from the compliance calibration method by using.20.21 25.283 I dC G --- IC- 2t da where t is the specimen thickness. Both P, and dClda can be accurately determined from the test results, as will be discussed in the Experimental Section. The compliance relationship with respect to the crack length, dClda, can also be expressed analytically by28 dC 3x2 da - G Wr3 +( t) -- where G is the elastic shear modulus of the material, x, W, and t are specimen dimensions shown in Fig. 1, and +(t) is the thickness correction factor, which may be expressed by28 $(t) = 1 - 0.6302(2t/W) + 1.20(2t1W)e-“’(2‘w (3) v2 t Fig. 1. Double torsion test configuration. The theoretical analysis of the DT technique assumes that the crack front is straight and orthogonal to the plane of the specimen. However, in most materials the crack front profile is significantly curved, with the furthest advance along the bot￾tom face of the specimen.24 Because of this, Evansz3 suggests that the compliance calibration be obtained experimentally rather than analytically. The consensus among most re~earchers~~.~~ has been that the crack velocity and crack length of interest is that of the leading edge of the crack, i.e., the furthest advanced bottom surface location. Tait, Fry, and Garrett33 have suggested that the optimum specimen dimensions are approximately as follows: width = W; length, L = 3W; and thickness = W16 to W/15. Greater thicknesses complicate the analysis because of the significant contact stresses generated by the interaction of the torsion arms (the two sides of the specimen separated by the crack). Another area of contention for the DT test configuration is the length of crack over which the stress intensity may be considered to be independent of crack length (constant-K re￾gion). Trantina30 showed that this assumption is valid for crack lengths greater than 0.55 W, and ligaments (uncracked portion) greater than 0.65W. Tait et found a consensus over several researchers which indicated that at least the central one-third to one-half of the specimen length (for UW = 3) should offer constant-K conditions. Shetty and Virkar34 stated that the valid region decreases as the specimen length-to-width ratio de￾creases. In. Experimental Procedure The materials used in this study were unidirectional ([O],,) Nicalon fiberlCAS-11 matrix composites. They were supplied by Coming, Inc. in the form of square plates, approximately 15 cm by 15 cm and about 3 mm thick. DT tests were conducted on an Instron Model 8562 servo￾electric universal testing machine. Tests were performed with a 500 lb LeBow load cell mounted below the standard 20 OOO lb Instron load cell. This lower capacity load cell was used be￾cause of its better resolution in the load range of interest. An Inconel rod was attached to the load cell for load application. A flat compression table was mounted to the bottom grip to hold the DT fixture base. A photograph of the loading section of the DT test system is shown in Fig. 2. The test fixture used for the DT test was similar to the one designed by Professor Leon Chuck from the University of Day￾ton.35 A close-up of the fixture, with a composite specimen mounted, appears in Fig. 3. The fore-and-aft, and side-to-side alignment of the base and the upper fixture is adjusted by modifying the length of the two pivot rollers. Once the speci￾men and fixture are aligned during setup, the support fixture for the top portion plays no role during loading. The fixture, the base plate, and the loading rod were all made from Inconel, a nickel-based superalloy, for use at temperatures up to 1000°C. The testing machine was equipped with a slot type, 2 Zone Model “D”, furnace, mounted on rails, supplied by Instron Corp. This furnace which was originally designed for tensile tests was enlarged axially by attaching alumina insulation plates to create a box-shaped furnace positioned over the slot furnace. Internal dimensions of the box were 15 cm x 15 cm. For those tests performed at high temperatures using the en￾larged furnace, the temperature was read by inserting a ther￾mocouple wire into the furnace next to the DT specimen. DT test specimen blanks of 2.18 cm x 6.70 cm were cut from the composite plates using a diamond saw and a diamond￾coated grinding wheel. Initial notches were cut into one side of the specimens using the 1 mm thick diamond saw. To deter￾mine the compliance vs crack length calibration, dClda, six to eight specimens were prepared and a different notch length was machined into each specimen. The introduced notch length ranged from 1.80 to 5.08 cm

1814 Journal of the American Ceramic Society-Kahraman et al Vol. 80. No. 7 a SLOPE 1 0.04 Displacement(E)mm Fig 4. Specimen compliance(C) calibration load vs displacement plot. 6.54cm 0.10cm Fig. 2. Loading section of the double torsion test system VIEW EDGE VIEW Flg. 5. Double torsion specimen for fracture testing, including notch Ig. 3. Photograph of the double torsion test fixture root detail The compliance calibration tests were conducted at a loading Before the dt tests to failure were performed the notch ti rate of 89 N/min, while plotting load vs piston displacement. were machined such that the thickness of the uncracked part of Specimens were loaded to low loads to avoid fracture during the specimen at the notch tip tapered from very thin to the full compliance calibration. Only the compliance of the specimen hickness as shown in Fig. 5. The top face of the specimen was before any appreciable crack growth occurred was of interest in the face with the longest notch, corresponding to the full thick the calibration. The sharpness of the crack tip was therefore not ness. This facilitated the initiation of a sharp precrack in the critical in determining dc/da specimen, as suggested by Tait et al. 33 Since the stress intensity A representative load(P)vs displacement(d) curve of a is inversely proportional to r,2325.3133 a reduction of the ompliance calibration test is shown in Fig. 4. Compliance, C, thickness by taper from full thickness to zero at the crack tip was determined by calculatin ing the slope of the linear portion of results in very high stress intensities at the P-d curve during loading. The nonlinearities in the lot facilitates the formation of a sharp crack at low loads. The load portion of the P-d curves may be attributed to the gradual crack can thus initiate at loads well below those required to accumulation of forces in the fixture. 5 The change in compli cause fast fracture of the full thickness. As noted by evans, ance with crack length, dClda, was calculated by performing a a sharp precrack is necessary for valid fracture toughness tests linear regression of compliance(C) vs crack length(a) in ceramics. The critical load values obtained from blunt cracks

1814 0.10 cm Journal of the American Ceramic Society--Kahraman et al. Vol. 80, No. 7 -6.69 c- -6.54 cm-----) Fig. 2. Loading section of the double torsion test system. Fig. 3. Photograph of the double torsion test fixture. The compliance calibration tests were conducted at a loading rate of 89 N/min, while plotting load vs piston displacement. Specimens were loaded to low loads to avoid fracture during compliance calibration. Only the compliance of the specimen before any appreciable crack growth occurred was of interest in the calibration. The sharpness of the crack tip was therefore not critical in determining dC/da. A representative load (P) vs displacement (d) curve of a compliance calibration test is shown in Fig. 4. Compliance, C, was determined by calculating the slope of the linear portion of the P-d curve during loading. The nonlinearities in the low￾load portion of the P-d curves may be attributed to the gradual accumulation of forces in the fixture.35 The change in compli￾ance with crack length, dC/du, was calculated by performing a linear regression of compliance (C) vs crack length (a). Displacement (6), rnm Fig. 4. Specimen compliance (0 calibration load vs displacement plot. Fig. 5. Double torsion specimen for fracture testing, including notch root detail. Before the DT tests to failure were performed, the notch tips were machined such that the thickness of the uncracked part of the specimen at the notch tip tapered from very thin to the full thickness as shown in Fig. 5. The top face of the specimen was the face with the longest notch, corresponding to the full thick￾ness. This facilitated the initiation of a sharp precrack in the specimen, as suggested by Tait et Since the stress intensity is inversely proportional to t1n,23.25*31*33 a reduction of the thickness by taper from full thickness to zero at the crack tip results in very high stress intensities at first loading which facilitates the formation of a sharp crack at low loads. The crack can thus initiate at loads well below those required to cause fast fracture of the full thickness. As noted by Evans,23 a sharp precrack is necessary for valid fracture toughness tests in ceramics. The critical load values obtained from blunt cracks

July 1997 High-Temperature Transverse Fracture Toughness a 100% increase in g tests to failure were performed with load-point displace- ment rates ranging from 0. 1 mm/h to 0.05 mm/min. a repre sentative P-d curve for the dt tests to failure is shown in Fig 6. Failure did not occur by a sudden total catastrophic fracture of the specimen, but rather by a gradual propagation of the E crack with a gradually increasing nonlinearity(decreasing ope). Fiber bridging might occur during the test due to the inevitable misalignment of fibers in the material, and there might be fibers which cross the plane of the crack.36-38As the crack opens, these fibers will be extracted from the matrix requiring additional applied load. Once crack propagation has become well established, fiber bridging will dominate the frac as it extends, in contrast to a fairly constant load for stable 6 ture energy. This results in an increasing load to grow the crack crack propagation in homogeneous materials. IHowever, the region of nonlinearity in Fig. 6 is small, suggesting that the b. i fibers are not very efficient at bridging the crack in the dt specimens(as will be discussed in the Results section), in contrast to the significant matrix crack interactions with the fibers reported for other geometries such as double cantilever beam(dCB)6.37 and flexure specimens. 38 Still the gradual failure made it difficult to precisely determine the critical load, In order to better define Pe, the offset procedure,35 was employed. P. was determined from the P-d curve with the Flg. 7. Compliance vs crack length for glass double torsion speci secant of 5% lower slope than the original elastic slope(slope mens. of the linear portion of the curve)as shown in Fig The transverse fracture toughness, GIc, was then calculated from Eq. (1) glass slides with known fracture toughness as a calibration of Fiber-matrix bond strength measurements were done utiliz. the Dt test apparatus and procedure. Experimental details are ing the microdebonding technique, where individual fibers are in Ref. 42 The compliance vs crack length plot and the slope(dC/da) obtained by lir is shown in Fig. 7. The mental dC/da, 0.000140 N-1, is in good agreement with the calculated dcyda(using 2), which is0000136N-1.4 I. Results Note that the compliance-axis-intercept is not zero for the experimental compliance calibration. This nonzero intercept is Before performing any fracture test on Nicalon/CAS-II com- attributed to the compliance of the uncracked part of the speci posite specimens, the reliability of the DT test results was men.28 The arms( two sides across the crack) of the dt s monstrated by the preliminary tests on soda-lime-silicate men are attached not to a ri ate, but to a plate(uncracked portion of the specimen) with some compliance which is inde pendent of crack length. The compliance of an uncracked specimen(a 0)is also included in Fig. 7. It falls close to the The fracture toughness, GIe, values were calculated from Eq appears from the data that the crack experiences constant-K conditions for about 0.1 <alL<0.7. This is the expected valid range from the literature.33 The Gc data in this constant-K region scatter in the range of 5.5-7.5 J/m2 with an average value of about 6.4 J/m2 Ke values in the literature are found to be in the range of 43.44 6.35-8.3 J/m2,42 Thus, the experimental data correlate well with those found in literature (1)Composite Compliance Calibration DT compliance calibration tests were performed on compos 54 Offset Compiance ite specimens with different premachined crack lengths. The resulting sample compliance vs crack length plot of those test -squares linear regress men compliance vs crack length data was done to determine the compliance calibration, dC/da, and the zero crack length inter cept. Note again that the compliance-axis -intercept is not zero for the experimental compliance calibration, as discussed 0,000020 060.080.10 above Displacement(5), mm The experimental dC/da, the slope of C-a plot in Fig. 9, is 5.35 x 10-6 N-. The analytical dc/da was calculated to cal load P than 2%

July 1997 High-Temperature Transverse Fracture Toughness 1815 from DT specimens may be high by as much as 40% causing about a 100% increase in The tests to failure were performed with load-point displace￾ment rates ranging from 0.1 mmfl.1 to 0.05 dmin. A repre￾sentative P-d curve for the DT tests to failure is shown in Fig. 6. Failure did not occur by a sudden total catastrophic fracture of the specimen, but rather by a gradual propagation of the crack with a gradually increasing nonlinearity (decreasing slope). Fiber bridging might occur during the test due to the inevitable misalignment of fibers in the material, and there might be fibers which cross the plane of the crack.3c38 As the crack opens, these fibers will be extracted from the matrix, requiring additional applied load. Once crack propagation has become well established, fiber bridging will dominate the frac￾ture energy. This results in an increasing load to grow the crack as it extends, in contrast to a fairly constant load for stable crack propagation in homogeneous materials.21 However, the region of nonlinearity in Fig. 6 is small, suggesting that the fibers are not very efficient at bridging the crack in the DT specimens (as will be discussed in the Results section), in contrast to the significant matrix crack interactions with the fibers reported for other geometries such as double cantilever beam (DCB)36,37 and flexure specimens.38 Still the gradual failure made it difficult to precisely determine the critical load, P,, at which crack propagation initiated. was employed. P, was determined from the P-d curve with the secant of 5% lower slope than the original elastic slope (slope of the linear portion of the curve) as shown in Fig. 6. The transverse fracture toughness, GI,, was then calculated from Eq. (1). Fiber-matrix bond strength measurements were done utiliz￾ing the microdebonding technique, where individual fibers are compressively loaded on a polished surface to produce debond￾ing.4"+2 In order to better define P,, the offset IV. Results Before performing any fracture test on NicalodCAS-II com￾posite specimens, the reliability of the DT test results was demonstrated by the preliminary tests on soda-limesilicate 250 200 Z 150 a U 0 h " 2 100 50 I I 0.00 002 0.04 0.06 0.08 0.10 0.12 0,14 Displacement (6), mm Fig. 6. Specimen load vs displacement plot from the double torsion tests to failure, including 5% offset compliance to determine the criti￾cal load, P,. 1 I I I I 0.00 0.0 1 0.02 0.03 0.04 0.05 0.06 Crack Lenglh (a), m Fig. 7. Compliance vs crack length for glass double torsion speci￾mens. glass slides with known fracture toughness as a calibration of the DT test apparatus and procedure. Experimental details are in Ref. 42. The compliance vs crack length plot and the slope (dC/du) obtained by linear regression is shown in Fig. 7. The experi￾mental dC/da, 0.000140 N-I, is in good agreement with the calculated dCYda (using Eq. (2)), which is 0.000136 N-1.42 Note that the compliance-axis-intercept is not zero for the experimental compliance calibration. This nonzero intercept is attributed to the compliance of the uncracked part of the speci￾men.28 The arms (two sides across the crack) of the DT speci￾men are attached not to a rigid plate, but to a plate (uncracked portion of the specimen) with some compliance which is inde￾pendent of crack length. The compliance of an uncracked specimen (a = 0) is also included in Fig. 7. It falls close to the trend-line intercept. The fracture toughness, GI,, values were calculated from Eq. (1) and plotted against crack length as shown in Fig. 8. It appears from the data that the crack experiences constant-K conditions for about 0.1 < a/L < 0.7. This is the expected valid range from the literature.33 The GI, data in this constant4 region scatter in the range of 5.5-7.5 J/m2 with an average value of about 6.4 J/mz. K,, values in the literature are found to be in the range of 0.7-0.8 MPa - m1'2,43,44 corresponding to a GI, range of about 6.35-8.3 J/m2.42 Thus, the experimental data correlate quite well with those found in literature. (1) Composite Compliance Calibration DT compliance calibration tests were performed on compos￾ite specimens with different premachined crack lengths. The resulting sample compliance vs crack length plot of those tests is shown in Fig. 9. A least-squares linear regression of speci￾men compliance vs crack length data was done to determine the Compliance calibration, dC/da, and the zero crack length inter￾cept. Note again that the compliance-axis-intercept is not zero for the experimental compliance calibration, as discussed above. The experimental dC/da, the slope of C-a plot in Fig. 9, is 5.35 x lod N-I. The analytical dC/da was calculated to be 5.45 x lo4 N-' using Eq. (2).42 The difference between the experimental and the analytical values of dC/da is less than 2%

I816 Joumal of the American Ceramic Society--Kahraman et al. Vol. 80. No. 7 Fig. 8. Critical strain energy release rate(GIe)measured on srack Flg. 9. Compliance vs crack length plot for Nicalon/CAS-lI double as a function of torsion specimens. men ness values for two Nicalon/CAS-lI composite plates are close (2) Room-Temperature Test Results to that of the CAS matrix, reported as 51.5 J/m2 at room tem Composite DT specimens were cut and notched as shown in perature. 7Schutz35 determined a lower Gis for Nicalon/CAS- Fig. 5. As stated earlier, the tapered notch facilitates the ini- Il, about 34 J/m2 with experimental compliance calibration tiation of a sharp precrack in the specimen at a load lower than (about 41 J/m2 if analytical compliance calibration was used) Pe, which is necessary for fracture toughness tests to be valid. (The difference between the analytical and experimental dc/d This was demonstrated by loading a specimen up to about 85% was about 20%, compared to 2% in this study of the critical load. The notch tip in the specimen was exam- Although matrix crack-fiber interactions, as the crack front ined under the optical microscope before and after the test. meanders and crosses over inclined fibers, are reported to While there was no change on the upper face of the specimen, dominate the measured fracture loads in conventional double up pictures of the notch tip on the bottom side of the specimen of this study the fibers do not appear to be very effective ts cantilever beam(DCB)specimens according to the res before the loading(without a sharp crack) and after the loading bridging the crack in the DT specimens. The bridging fibers in (with the introduced sharp crack)are shown in Figs. 10 and 11 o.The sharp precrack extended about 7 mm on the bottom face pressure is achieved at a crack opening of about 60 um. That crack(notch) length on the top face. This observation was in (see next section). Thus, there might not be enough crack open good agreement with the curved front profile discussed earlier ing in the DT specimen to show a significant toughness in- The fracture toughness test data obtained in this study therefore crease due to the fibers as compared to the DCB specimen results from naturally developed sharp cracks prior to failure The relationship between the transverse fracture toughne ()High-Temperature Transverse Fracture Results Gre, and the initial crack length, a, for Nicalon/CAS-lI DT The same analytical dC/da which was determined at room specimens from two plates is illustrated in Fig. 12. The crack mperature was used for high-temperature transverse fracture length used in the plot is the approximate length of the crack on toughness calculations, assuming that the composite retains its the bottom side of the specimen after the initiation of the sharp room-temperature compliance at temperatures up to 1000C crack at the notch tip. It was demonstrated above that after the This was demonstrated by determining the compliance of a initiation of a sharp crack at the notch tip at a load lower than Nicalon/CAS-IIDT specimen both at room temperature and at Pe the crack length on the bottom face becomes about 4 mm 1000oC. The change in compliance between room temperature longer than the notch length on the top face nd 1000C was small. less than 5%. This is consistent with the There seems to be no discernible trend between the lack of any significant decrease in modulus observed over this sured transverse fracture toughness and the initial crack length range. 5 The specimens tested at different temperatures had the in the range of crack lengths used. This confirms the con same(top face)crack length of about 3. 28 cm stant-K characteristics of this specimen type, as already dem- A transverse fracture toughness(Gre)vs time to failure(test- onstrated on glass specimens for about 0.1 alL <0.7. The ing time)plot of Nicalon/CAS-Il specimens for room tempera- data point for alL =0.73 may not be valid because of the long ture, 800 C and 1000@C is shown in Fig 13. There seems to b initial crack length, which might be out of the constant-k no significant effect of test rate on transverse fracture tough Considering only the specimens with 0.1 a/L <0.7, the 38-61 J/m with an average value of 48 J/m2. On the or ness at room temperature. The Gre data scatter in the ra average transverse fracture toughness value is 49 and 50 J/m hand, the transverse fracture toughness of Nicalon/CASIl for plates 1 and 3, respectively. These measured fracture tough- composite was determined to be lower than 38 J/m- for the

1816 Journal of the American Ceramic Society-Kahraman et al. Vol. 80. No. 7 f f \ I I I I 0.2 0.4 0.6 0.8 a/l Fig. 8. Critical strain energy release rate (G,c) measured on soda￾lime-silicate glass double torsion specimens as a function of crack length-to-specimen length ratio (dL). (2) Room-Temperature Test Results Composite DT specimens were cut and notched as shown in Fig. 5. As stated earlier, the tapered notch facilitates the ini￾tiation of a sharp precrack in the specimen at a load lower than P,, which is necessary for fracture toughness tests to be valid. This was demonstrated by loading a specimen up to about 85% of the critical load. The notch tip in the specimen was exam￾ined under the optical microscope before and after the test. While there was no change on the upper face of the specimen, a sharp crack initiated at the notch tip on the lower face. Close￾up pictures of the notch tip on the bottom side of the specimen before the loading (without a sharp crack) and after the loading (with the introduced sharp crack) are shown in Figs. 10 and 11, respectively. The sharp precrack extended about 7 mm on the bottom face making the crack length on this face 4 mm longer than the crack (notch) length on the top face. This observation was in good agreement with the curved front profile discussed earlier. The fracture toughness test data obtained in this study therefore results from naturally developed sharp cracks prior to failure. The relationship between the transverse fracture toughness, G,, and the initial crack length, a, for NicalodCAS-11 DT specimens from two plates is illustrated in Fig. 12. The crack length used in the plot is the approximate length of the crack on the bottom side of the specimen after the initiation of the sharp crack at the notch tip. It was demonstrated above that after the initiation of a sharp crack at the notch tip at a load lower than P, the crack length on the bottom face becomes about 4 mm longer than the notch length on the top face. There seems to be no discernible trend between the mea￾sured transverse fracture toughness and the initial crack length in the range of crack lengths used. This confirms the con￾stant-K characteristics of this specimen type, as already dem￾onstrated on glass specimens for about 0.1 < a/L < 0.7. The data point for a/L = 0.73 may not be valid because of the long initial crack length, which might be out of the constant-K range. Considering only the specimens with 0.1 < a/L < 0.7, the average transverse fracture toughness value is 49 and 50 J/m2 for plates 1 and 3, respectively. These measured fracture tough- 0 0.00 L 0.01 0.02 0.03 0 Crack Length (a). m Fig. 9. Compliance vs crack length plot for NicalodCAS-I1 double torsion specimens. ness values for two NicalodCAS-I1 composite plates are close to that of the CAS matrix, reported as 51.5 J/m2 at room tem￾perature." Sch~tz~~ determined a lower G, for NicalodCAS￾II, about 34 J/m2 with experimental compliance calibration (about 41 J/m2 if analytical compliance calibration was used). (The difference between the analytical and experimental dC/da was about 20%. compared to 2% in this study.) Although matrix crack-fiber interactions, as the crack front meanders and crosses over inclined fibers, are reported to dominate the measured fracture loads in conventional double cantilever beam (DCB) specimen^,^^^* according to the results of this study the fibers do not appear to be very effective at bridging the crack in the DT specimens. The bridging fibers in the crack wake exert a crack closing pressure for a given crack opening and according to Kaute et a1.36 the maximum closing pressure is achieved at a crack opening of about 60 pm. That large of a crack opening is never achieved in DT specimens (see next section). Thus, there might not be enough crack open￾ing in the DT specimen to show a significant toughness in￾crease due to the fibers as compared to the DCB specimen. (3) High-Temperature Transverse Fracture Results The same analytical dC/da which was determined at room temperature was used for high-temperature transverse fracture toughness calculations, assuming that the composite retains its room-temperature compliance at temperatures up to 1O0O"C. This was demonstrated by determining the compliance of a NicalodCAS-11 DT specimen both at room temperature and at 1OOO"C. The change in compliance between room temperature and 1OOO"C was small, less than 5%. This is consistent with the lack of any significant decrease in modulus observed over this range.35 The specimens tested at different temperatures had the same (top face) crack length of about 3.28 cm. A transverse fracture toughness (GI,) vs time to failure (test￾ing time) plot of NicalodCAS-I1 specimens for room tempera￾ture, 800°C and 1000°C is shown in Fig. 13. There seems to be no significant effect of test rate on transverse fracture tough￾ness at room temperature. The GI, data scatter in the range of 38-61 J/m2 with an average value of 48 J/m2. On the other hand, the transverse fracture toughness of NicalodCAS-I1 composite was determined to be lower than 38 J/mz for the

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1818 Joumal of the American Ceramic Sociery-Kahraman et al. Vol. 80. No. 7 Fig. 12. Room-temperature fracture toughness vs crack length-to- specimen length ratio for Nicalon/CAs-Il double torsion specimens. Time to failure, h Transverse fracture toughness vs time to failure plot of Nicalon/CAS-Il compos fracture tests performed at 800 and 1000C. Also, the test rate had some effect on the composite toughness at 1000oC. The average Gre measured on the samples fractured in 9-10 h was about 20% lower than that of the ones fractured in less than 2 h. Overall, there is a reduction in Gle from about 48 J/m2down to under 38 J/m2 as the temperature increases from room tem- perature to 800 or 1000oC. It was established in a related study-2 that cracks grow in the of interphase oxidation in the interior of the co r,permeation of oxygen through the crack, after it seems to cause the oxidation of the interphase, in- the fiber-matrix bond strength. The crack then opens

1818 Journal of the American Ceramic Society-Kahraman et al. 9 30- c) 20 - ID￾Vol. 80, No. 7 0 0.4 0.5 0.6 0.7 c 4 A plate 1 plate 3 Fig. 12. Room-temperature fracture toughness vs crack length-to￾specimen length ratio for NicalodCAS-II double torsion specimens. 60/ 50 t 0 0 2 4 6 8 10 Time to Failure, h t 20°C 800'T A lOOO'JC 3 Fig. 13. Transverse fracture toughness vs time to failure plot of NicalodCAS-II composite. fracture tests performed at 800" and 1OOO"C. Also, the test rate had some effect on the composite toughness at 1OOO"C. The average G, measured on the samples fractured in 9-10 h was about 20% lower than that of the ones fractured in less than 2 h. Overall, there is a reduction in GI, from about 48 J/mz down to under 38 J/m2 as the temperature increases from room tem￾perature to 800" or 1OOO"C. that cracks grow in the absence of interphase oxidation in the interior of the composite. However, permeation of oxygen through the crack, after it initiates, seems to cause the oxidation of the interphase, in￾creasing the fiber-matrix bond strength. The crack then opens, It was established in a related 9 8 00 * m 8 E .I B c .I v) I. s m % c '43 8 v) 2 U

July 1997 High-Temperature Transverse Fracture Toughness 1819 breaking the fibers bridging the crack surface without pulling from the micrograph of the crack in Fig. 14. The crack is tightly them out of the matrix. This process is test time dependent, i.e., closed near the top side. Crack opening near the top side is no e slower the fracture the more effective the oxidation process more than a few micrometers even after handling between The increased interfacial bond strength results in a planar frac- test and the mounting of the sample(which was cut from the ture surface without noticeable fiber pullout, possibly lowering cracked DT specimen). The crack opening, which was prob- e transverse fracture toughness of the composite. However, bly much smaller than that during and after the failure of the the decrease in the transverse composite fracture toughness at DT specimen, might not have allowed air to penetrate through 800 and 1000@C is believed to be caused mainly by lower to the top side matrix fracture toughness at high temperatures relative to that The possible occurrence of high-temperature stress corrosion at room temperature, As much as 50% reduction has been crack growth was also explored by performing two dt tests observed in CAS matrix toughness at 1000 C 45 holding the load(P)constant to give Gr of about 20 and 25 J/m Microdebonding tests were also performed along the crack for about 3 h at 800C(G, is determined from Eq (I)replacing n cross sections of dt specimens fractured at temperatures G, and p。 by gr and p, ely). Examination of the from 20 to 1000C(1-2 h long tests ) One such cross section tested specimens under the light microscope did not reveal any fibers bridging the crack surface crack growth Cracks grow only when Gy is close to the critical aused the two arms of the dt specimens to remain intact even value, GIc, indicating that no significant stress corrosion crack after the failure. The crack is more opened on the bottom side growth parallel to the fibers is observed in the Nicalon/CAS-II of the specimen(with a maximum crack opening of about 30 composite at lower GI, at least for the experiments carried out um)and less opened on the top side due to the nature of the in this study. However, Spearing et al. 46 reports significant loading as discussed previously. Only the yery close stress corrosion matrix cracking on the same material system (within less than a fiber diameter distance)to the crack were under longitudinal loading conditions, at str required to develop matrix cracks in short-duration, monotonic Fiber-matrix bond strength distributions along the crack be. loading tests, even at room temperature but for much longer tween the two(top and bottom) surfaces of the DT specimens stress application times(as long as 278 h) Some environmental stress cracking might be expected by in Fig. 15. No significant change in fiber-matrix bond strengt oxidative removal of the interfacial carbon layer at high tem- was observed from 20 to 600C. The most significant increase peratures even for the short exposure times of this study. How- in the interfacial strength was observed at 1000C. However, a ever, this seems to be prevented by subsequent replacement of critical observation is that oxidation was effective only halfway the oxidized carbon layer with a new interfacial phase closing across the fracture surface from the bottom side while the bone the gap between fiber and matrix, and sealing the oxidation strengths on the other half were similar to those of control process off. 14 4174 opposites(except for the oxidized fibers very close to the top surface exposed to the environment). The results for the frac ture at 800%C were similar in nature but with lower increases in V. Summary and Conclusions Dt test technique was successfully utilized to determine The explanation for this selective oxidation may be drawn fracture toughness(Grc) for cracks parallel to the fibers in a c10 top bottom Dislance from Top Surfoce/Thickness fig. 15. Interfacial bond strength distribution along the crack from top to bottom face for DT specimens fractured(1-2 h long tests)at different

July 1997 High-Temperature Transverse Fracture Toughness 1819 breaking the fibers bridging the crack surface without pulling them out of the matrix. This process is test time dependent, i.e., the slower the fracture the more effective the oxidation process. The increased interfacial bond strength results in a planar frac￾ture surface without noticeable fiber pullout, possibly lowering the transverse fracture toughness of the composite. However, the decrease in the transverse composite fracture toughness at 800" and 1OOO"C is believed to be caused mainly by lower matrix fracture toughness at high temperatures relative to that at room temperature. As much as 50% reduction has been observed in CAS matrix toughness at 1000°C."5 Microdebonding tests were also performed along the crack on cross sections of DT specimens fractured at temperatures from 20" to 1OOO"C (1-2 h long tests). One such cross section is shown in Fig. 14. The fibers bridging the crack surface caused the two arms of the DT specimens to remain intact even after the failure. The crack is more opened on the bottom side of the specimen (with a maximum crack opening of about 30 km) and less opened on the top side due to the nature of the loading as discussed previously. Only the fibers very close (within less than a fiber diameter distance) to the crack were tested. Fiber-matrix bond strength distributions along the crack be￾tween the two (top and bottom) surfaces of the DT specimens tested at temperatures of 20", 600", 800", and 1OOO"C are given in Fig. 15. No significant change in fiber-matrix bond strength was observed from 20" to 600°C. The most significant increase in the interfacial strength was observed at 1OOO"C. However, a critical observation is that oxidation was effective only halfway across the fracture surface from the bottom side, while the bond strengths on the other half were similar to those of control composites (except for the oxidized fibers very close to the top surface exposed to the environment). The results for the frac￾ture at 800°C were similar in nature but with lower increases in bond strength. The explanation for this selective oxidation may be drawn from the micrograph of the crack in Fig. 14. The crack is tightly closed near the top side. Crack opening near the top side is not more than a few micrometers even after handling between the test and the mounting of the sample (which was cut from the cracked DT specimen). The crack opening, which was prob￾ably much smaller than that during and after the failure of the DT specimen, might not have allowed air to penetrate through to the top side. The possible occurrence of high-temperature stress corrosion crack growth was also explored by performing two DT tests holding the load (P) constant to give GI of about 20 and 25 J/mz for about 3 h at 800°C (GI is determined from Eq. (1) replacing GI, and P, by GI and P, respectively). Examination of the tested specimens under the light microscope did not reveal any crack growth. Cracks grow only when GI is close to the critical value, GI,, indicating that no significant stress corrosion crack growth parallel to the fibers is observed in the NicalodCAS-Il composite at lower GI, at least for the experiments carried out in this study. However, Spearing et aL46 reports significant stress corrosion matrix cracking on the same material system under longitudinal loading conditions, at stresses below that required to develop matrix cracks in short-duration, monotonic loading tests, even at room temperature but for much longer stress application times (as long as 278 h). Some environmental stress cracking might be expected by oxidative removal of the interfacial carbon layer at high tem￾peratures even for the short exposure times of this study. How￾ever, this seems to be prevented by subsequent replacement of the oxidized carbon layer with a new interfacial phase closing the gap between fiber and matrix, and sealing the oxidation process 0ff.14-17*42 V. Summary and Conclusions DT test technique was successfully utilized to determine fracture toughness (GI,) for cracks parallel to the fibers in a 0 a 2 x * 20oc 0 600OC 8OOOC A lO0O~C 0 0.0 0.2 A 4A 4 1. rn 0. 0 A x I T 1 I . 0. bottom I I I I 0.4 0.6 0.8 I Dislance from Tap Surface/Thickness 0 Fig. 15. temperatures (bond strengths with mows exceeded the measurement capacity of the apparatus). Interfacial bond strength distribution along the crack from top to bottom face for DT specimens fractured (1-2 h long tests) at different

820 Journal of the American Ceramic Sociery-Kahraman ef al Vol. 80. No. 7 ceramic matrix composite. The data were obtained for Nicalon/ ings on Interfaces in Composites 193-204. Edited by C. G. Pantano and CAS-Il over the range from 20 to 1000oC. GIe did decrease urgh, PA, 1990. moderately with increasing temperature(as does the bulk ma Evaluating Trans- trix), but no evidence of an interphase oxidizing effect on crack er.Scl.Let,1u12]51l-55(1982) growth could be found. Even after fracturing and being ex- posedonthefracturesurfaceduringcool-down,interfacesnearF2lugsmodeglaminationandTRansverseCracking.compos,.Mater20 bond strength. The more open side showed higher bond Publishers, Dordrecht, Nether strength(and a flatter fracture surface42 as expected ). Cracks would not grow in the oxidizing environment at GI values Bradt, D P H New York, 1973 htly below Grc(for a 3 h exposure time), and oxidation did ng Time-Dependent Failure Charact not occur on the part of the fracture surface which was cracked Mater. Sci. 7(10) 1137-45(1972) but not widely opened during the test 2D P. williams and A G. Evans, ""Simple Method for Studying Slow Crack Acknowle ch would not have occurred without the 26A. G. Evans and S M. Wiederhorm, ""Crack Propagation and Failu tride at Elevated Te nd University of Petroleum an onle enck ies arad wr e: p. 4 3-9 edited byg l Peat chap mat and an New York. 1969. E.R. Fuller, Jr, "Fracture Mechanics Applied to Brittle Materials: References P man. American Society for Testing and Materials, Philadelphia, PA, 1979 1-56 in High Temperature/High Performance Composites, Edited 20. Sano,"'A Revision of the Double Torsion Technique for Brittle Mate Lemkey, S.G. Fishman, A G. Evans, and J.R. Strife. Materials Re- rials, J. Mater. Sci, 23, 2505-1l(i988) J O. Stiegler, "Structural Ceramics R& D, " Adv. Mater. Ceram Soc., 60 [7-8]33841(197 Processes,138355-61(1990 3C. G. Annis and J S. Cargill, Fracture Mechanics of Ceramics: pp. 737 ties for Com- 44. Edited C. Bradt, D. P H. Hasselman, and F. F. Lange. Plenum Press, New York, 1973 3P. S. Leevers, Crack Front Shape Effects in the Double Torsion Test, J. R B. Tait, P. R. Fry, and G.G. Garrett, " Review and Evaluation of the with High Tensile Strength, J. Am. Ceram. Soc., 59[7-8] Double Torsion Technique for Fracture Toughness and Fatigue Testing 324-27(1976) "T- L. Mah, M. G. Mendiratta, A.P. Katz, and K S. Mazdiyasni,"Recent D. K Shetty and A. v. virkar, mination of useful ri Specimens, J Am. Ceram Soc., 61[1-2]93-94 Am. Ceran Soc. Bull, 66 (2]304-18(1987 of Fiber matrix 33J. B Schutz, "Test Methods and Analysis for Glass-Ceramic Matrix Com nic ma- triNicalon Fiber Composites"; Pp, 349-60 in Proceedings of the Ce R. Shercliff, and M. F. Ashby, ""Delamination, Fiber Composite Ceramics. Edited by r. e. Tressler of Ceramic Matrix Composites, Acta Metall, G L. Messing, C G. Pantano, and R. E. Newnham. Plenum Press, New York, 417]195970(1993). D. A. W. Kaute, H. R. Shercliff, and M. F. Ashby, 9R. I Kerans, R S. Hay, N I. Pagano, andT. A Parthasarathy "The Role of m. Ceram. Soc. BulL. 68 [2]42942(1989) Reinforced Brittle Matrix Composites, J. sci.29,3857-%6( E Y Luh and A G. Evans, "High Temperature Me Transverse 啊Pmm.D Modified R. L Stewart, K. Chyung, M. P. Taylor, and R F. Cooper, "Fracture of pp.87-108 in ASTM Special Te Publica- mics of Ceramics. Edited by r c. Bradt, D H. Grande, J F. Mandell, and K CC. Hong, Fiber-Matrix Bond M. Huger, D. Fargeot, and C. Gault, "Ultrasonic Characterization of chanisms in Nicalon/C/SiC Composites, "J. Am. Ceram. Soc., 77[10] sci,23,311-28(1988) 2554-60(1994) 42R. Kahraman, "" Influence of Fiber-Matrix Interphase on High Temperature "Oxidation Mechanisms and Ki Ph.D. the 3J. F. Shackelford, Introduction to Materials Science for rs. Mac uzzi and R. naslain, *Oxidation Mechanisms and Kinetics of ID- g Development and Failure of Fiber SCorning Inc, private communication to Dr. J. F. Mandell, 1992 einforced Ceramic matrix C Ph. D. Thesis, Massachusetts Institute A. G. Evans, ' Stress nidirectional Ceramic Matrix Composite, "J. Am. Ceram, Soc.. 77 R. A. Shimansky, H. T. Hahn, and N J. Salamon, ""Symposium Proceed- 0(1994

1820 Journal of the American Ceramic SocietpKahraman et al. Vol. 80, No. 7 ceramic matrix composite. The data were obtained for Nicalod CAS-II over the range from 20” to 1ooo”C. GI, did decrease moderately with increasing temperature (as does the bulk ma￾trix), but no evidence of an interphase oxidizing effect on crack growth could be found. Even after fracturing and being ex￾posed on the fracture surface during cool-down, interfaces near the less open side of the fracture surface showed no increase in bond strength. The more open side showed higher bond strength (and a flatter fracture surface4* as expected). Cracks would not grow in the oxidizing environment at GI values slightly below GI, (for a 3 h exposure time), and oxidation did not occur on the part of the fracture surface which was cracked but not widely opened during the test. Acknowledgments: This research would not have occurred without the composite materials supplied by Corning. Inc., their help is acknowledged. 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