ournal JAm. Ceram.So,6间6981-9002003) Characterization and High-Temperature Mechanical Behavior of an Oxide/Oxide Composite Larry P. Zawada, Randall S Hay, Shin S Lee, and James Staehler Materials and Manufacturing Directorate, Air Force Research Laboratory, Wright-Patterson Air Force Base, Ohio45433-7817 An oxide/oxide ceramie fiber-matrix composite( CMC) has porous, cracked matrices for high-temperature aerospace applica- structural applications. This CMC is called GEN-IV, and it and an aluminosilicate matrix is called GEN-IV. 4, 5,2 and it is has a porous and cracked aluminosilicate matrix reinforced by referenced in the following text as N610/AS. A review of the 3M Nextel 610 alumina fibers woven in a balanced eight mechanisms and mechanical properties of porous-matrix CMCs harness weave (SHSW). This CMC has been specifically has been given elsewhere. designed without an interphase between the fiber and matrix, The following characteristics of N610/AS porous-matrix com- and it relies on the porous matrix for flaw tolerance. Stress- posites are evaluated:()long-time phase and microstructural strain response is nearly linear to failure and without a stability; (i high-temperature, short-time stress-strain response, well-defined proportional limit in tension and compression. (iii) fatigue; and(iv)creep and creep rupture. Comparisons with In-plane shear and interlaminar strength increases with in- other CMCs are made and possible explanations for the composite easing temperature. The 1000C fatigue limit in air at 105 veles is 160 MPa and the residual tensile strength of run-out specimens is not affected by the fatigue loading. The creep- rupture resistance above 1000C is relatively poor, but it can Il. Materials and Experiments be improved with a more-creep-resistant fiber Composite Fabrication 3M Nextel 610 fibers are used to manufacture tile of N6lO/AS The Nextel 610 fiber is 99 wt% polycrystalline a-Al,O3, with a L. Introduction density of 3.88 g/cm, an average grain size of 0. 1 um, and an IGH fracture toughness and damage tolerance is engineered average filament diameter of -12 um. Eight harness satin weave (8HS W) cloth of Nextel 610 was prepregged with a mixture of fine (CMCs)by tailoring properties of the fiber-matrix interface. The Al2 O3 powder and a SiO2-forming polymer 14, I5 Twelve individ- ual prepregged cloths were stacked on top of each other as a fiber-matrix interface must deflect matrix cracks and allow fiber laminate. The laminate was warm molded in an autoclave to pullout afterward. 2 Mechanical properties of CMCs break down if the coating is not stable in the application environment.- produce a dense green-state ceramic tile. The tile was pressureless Carbon- and BN- are the usual fiber-matrix interphases sintered in air at -1000oC. This process removed organics SiC-fiber CMCs. Unfortunately, carbon coatings begin to oxidize converted the polymer to porous Sio at "450C, and the gap left by oxidation may fill with the siO2 oxidation product of SiC and form a strong fiber-matrix bond that (2) Experiments seriously degrades CMC mechanical properties. -> 8 Amorphous (A) Microstructure Characteriation: The composite den- nd imperfectly crystalline BN are moisture sensitive and easily sity (seven specimens) was measured using the Archimedes oxidize. Similar to SiO2, B2O3 forms a strong fiber-matrix principle and a helium pycnometer. Because the material readily bond that degrades CMC properties. Oxidation is a serious dsorbs moisture from the air, the specimens were carefull obstacle to long-term use of CMCs with carbon or BN fiber- outgassed before measurement. The total pore surface area was matrix interfaces at intermediate and high temperatures measured using the Brunauer-Emmett-Teller (BET) method An approach to flaw-tolerant CMCs that are also oxidation (eight different measurements on seven specimens). Fiber volume sistant is oxide/oxide CMCs with fibers that are " strongly Tractions were measured in three specimens polished at 45. to the bonded to a matrix deliberately made weak by incorporation of fiber axes. A microscope(Model Metallovert, Leitz)mounted with high porosity and microcracks. Instead of crack deflection and a video camera, video monitor, and computer running image sliding at fiber-matrix interfaces, fracture energy is dissipated by analyzer software(CUE-4, version 3.2, Olympus) was used for diffuse microcracking in the porous matrix. Early modeling of image analysis of the fiber volume fractions. these materials suggests that the fiber bundles must be hetere Fracture surfaces of failed specimens were characterized using enously distributed in the matrix, have higher coefficient of SEM. Creep rupture specimens were characterized using TEM thermal expansion (CTE) than the matrix, and have a Mode As-fabricated specimens and specimens heat-treated for 3000 h at fracture energy twice the Mode Il fracture energy of the ma- 982C(1800.F) were characterized using optical microscopy trix.9,20 General Electric has developed oxide/oxide CMCs with SEM(Model 360FE, Leica), and TEM operating at 200 kV(Model 2000 FX, JEOL, Tokyo, Japan). SEM specimens and TEM thin ections were prepared using a method described elsewhere. TEM specimens were mounted on copper grids and ion milled at F. Zok--contributing editor 7kV, with the last 15 min at 4 kV. Most specimens were carbon oated. Some of the TEM specimens were observed without carbon o that the fine structure of the matrix porosity resolved. For these specimens, care was taken to Manuscript N. 188004. Received January 17, 2002; approved January 28, 2003. mate the microscope for the particular condenser lens
Characterization and High-Temperature Mechanical Behavior of an Oxide/Oxide Composite Larry P. Zawada,* Randall S. Hay,* Shin S. Lee,* and James Staehler* Materials and Manufacturing Directorate, Air Force Research Laboratory, Wright-Patterson Air Force Base, Ohio 45433–7817 An oxide/oxide ceramic fiber–matrix composite (CMC) has been extensively characterized for high-temperature aerospace structural applications. This CMC is called GEN-IVTM, and it has a porous and cracked aluminosilicate matrix reinforced by 3M Nextel 610TM alumina fibers woven in a balanced eight harness weave (8HSW). This CMC has been specifically designed without an interphase between the fiber and matrix, and it relies on the porous matrix for flaw tolerance. Stress– strain response is nearly linear to failure and without a well-defined proportional limit in tension and compression. In-plane shear and interlaminar strength increases with increasing temperature. The 1000°C fatigue limit in air at 105 cycles is 160 MPa, and the residual tensile strength of run-out specimens is not affected by the fatigue loading. The creep– rupture resistance above 1000°C is relatively poor, but it can be improved with a more-creep-resistant fiber. I. Introduction HIGH fracture toughness and damage tolerance is engineered into most fiber-reinforced ceramic-matrix composites (CMCs) by tailoring properties of the fiber–matrix interface. The fiber–matrix interface must deflect matrix cracks and allow fiber pullout afterward.1,2 Mechanical properties of CMCs break down if the coating is not stable in the application environment.3–5 Carbon6–8 and BN9–11 are the usual fiber–matrix interphases in SiC-fiber CMCs. Unfortunately, carbon coatings begin to oxidize at 450°C, and the gap left by oxidation may fill with the SiO2 oxidation product of SiC and form a strong fiber–matrix bond that seriously degrades CMC mechanical properties.3–5,8 Amorphous and imperfectly crystalline BN are moisture sensitive and easily oxidize.10,12 Similar to SiO2, B2O3 forms a strong fiber–matrix bond that degrades CMC properties. Oxidation is a serious obstacle to long-term use of CMCs with carbon or BN fiber– matrix interfaces at intermediate and high temperatures. An approach to flaw-tolerant CMCs that are also oxidation resistant is oxide/oxide CMCs with fibers that are “strongly” bonded to a matrix deliberately made weak by incorporation of high porosity and microcracks.13–20 Instead of crack deflection and sliding at fiber–matrix interfaces, fracture energy is dissipated by diffuse microcracking in the porous matrix. Early modeling of these materials suggests that the fiber bundles must be heterogenously distributed in the matrix, have higher coefficient of thermal expansion (CTE) than the matrix, and have a Mode I fracture energy twice the Mode II fracture energy of the matrix.19,20 General Electric has developed oxide/oxide CMCs with porous, cracked matrices for high-temperature aerospace applications. The General Electric CMC with the 3M Nextel 610TM fiber and an aluminosilicate matrix is called GEN-IV,14,15,21 and it is referenced in the following text as N610/AS. A review of the mechanisms and mechanical properties of porous-matrix CMCs has been given elsewhere.22 The following characteristics of N610/AS porous-matrix composites are evaluated: (i) long-time phase and microstructural stability; (ii) high-temperature, short-time stress–strain response; (iii) fatigue; and (iv) creep and creep rupture. Comparisons with other CMCs are made and possible explanations for the composite behavior are discussed. II. Materials and Experiments (1) Composite Fabrication 3M Nextel 610 fibers are used to manufacture tile of N610/AS. The Nextel 610 fiber is 99 wt% polycrystalline -Al2O3, with a density of 3.88 g/cm3 , an average grain size of 0.1 m, and an average filament diameter of 12 m. Eight harness satin weave (8HSW) cloth of Nextel 610 was prepregged with a mixture of fine Al2O3 powder and a SiO2-forming polymer.14,15 Twelve individual prepregged cloths were stacked on top of each other as a laminate. The laminate was warm molded in an autoclave to produce a dense green-state ceramic tile. The tile was pressureless sintered in air at 1000°C. This process removed organics and converted the polymer to porous SiO2. (2) Experiments (A) Microstructure Characterization: The composite density (seven specimens) was measured using the Archimedes principle and a helium pycnometer. Because the material readily adsorbs moisture from the air, the specimens were carefully outgassed before measurement. The total pore surface area was measured using the Brunauer–Emmett–Teller (BET) method (eight different measurements on seven specimens). Fiber volume fractions were measured in three specimens polished at 45° to the fiber axes. A microscope (Model Metallovert, Leitz) mounted with a video camera, video monitor, and computer running image analyzer software (CUE-4, version 3.2, Olympus) was used for image analysis of the fiber volume fractions. Fracture surfaces of failed specimens were characterized using SEM. Creep rupture specimens were characterized using TEM. As-fabricated specimens and specimens heat-treated for 3000 h at 982°C (1800°F) were characterized using optical microscopy, SEM (Model 360FE, Leica), and TEM operating at 200 kV (Model 2000 FX, JEOL, Tokyo, Japan). SEM specimens and TEM thin sections were prepared using a method described elsewhere.23 TEM specimens were mounted on copper grids and ion milled at 7 kV, with the last 15 min at 4 kV. Most specimens were carbon coated. Some of the TEM specimens were observed without carbon coating so that the fine structure of the matrix porosity could be better resolved. For these specimens, care was taken to align and stigmate the microscope for the particular condenser lens setting used. F. Zok—contributing editor Manuscript No. 188004. Received January 17, 2002; approved January 28, 2003. *Member, American Ceramic Society. J. Am. Ceram. Soc., 86 [6] 981–90 (2003) 981 journal
982 Journal of the American Ceramic Sociery-Zawada et al. Vol 86. No 6 B) Mechanical Test Apparatus: A horizontal servohydraulic porous SiO,. A micrograph taken 45% to the warp and fill fibers is machine with rigid hydraulic clamping grips and quartz-lamp heating shown in Fig. 1. Parallel arrays of cracks in the matrix were was used for the tension, in-plane shear, creep rupture, and fatigue perpendicular to the cloth layers. The crack spacing was wider in tests. Test control, data acquisition, and interactive data analysis was the matrix-rich regions and smaller within the fiber tows. Presun done using the MATE program" on an IBM-compatible personal ably, these cracks formed by matrix shrinkage during sintering that computer(PC)linked to the test frame by an analog-to-digital board. was constrained by the cloth layers, similar to that observed for Temperature was measured using five S-type thermocouples bonded constrained sintering of films- and around inclusions.2--29These to each specimen with an alumina-based ceramic adhesive. a detailed cracks formed under relatively low stresses that could not exceed description of the test equipment was given elsewhere. The inter- twice the sintering stress. The fiber-matrix interfaces and the laminar strength and in-plane shear tests were done using a standard fibers themselves were not cracked vertical servohydraulic test machine with a box fumace that used The average skeletal densities of seven specimens were 3.62 igniter elements gcm using the Archimedes method and 3.65 g/cm'using the Monotonic tests consisted of testing two to three test specimens pycnometer, with a standard deviation of +0.06. An average bulk per condition. Fatigue and creep rupture testing consisted of testing density of 2.90 g/cm' was calculated from the immersion meas- only one test specimen at each stress level investigated. It was urements. The average fiber volume fraction was -30.7%+ recognized by the authors that this was an extremely limited set of 2.29%. The SiO, and Al,O3 volume fractions of the matrix were data on which to make scientific observations. However, the not measured, but the expected proportions following processing bjective was to explore the boundaries of mechanical behavior for were 87 wt%a-Al2O3 and 13 wt% amorphous SiO,(2. 2 g/cm) oxide/oxide CMCs and to determine if they warranted a more Therefore, the interconnected porosity in the entire composite was rigorous investigation for use in aerospace applications. -24%, or 35% in the matrix alone. Al2O3 occupied 51 vol% of the matrix and SiO2 14%. Because this matrix porosity was concen- were performed using stroke control with a 0.05 mm/s displace- trated in the SiO2, initial SiO, porosity was 71%. However, ment rate. Tension tests were performed using dogbone test sintering shrinkage cracks accounted for some of this porosity,the remainder was finely distributed in the Sio, that cemented the specimens. Tension tests were also done at 1000 and 1100 C in Al2 O3 grains together. The high specific surface area measured by with an -l.5 cm hot zone In high-temperature tests, each specimen was ramped to the test 3.98 mlg) was consistent with a large amount of fine intercon- temperature in 15 min and then equilibrated for - 20 min; the nected porosity. In contrast to N610/AS, dense glass-ceramic stress then was ramped up until the specimen failed. For in-plane composites of Nicalon/MAS were found to have a specific surface shear measurement, tension tests were performed using test spec area of <0.3 m/g imens with +45 fiber orientations. Residual room-temperature The mixture of a-AL,O3 and amorphous SiO, that forms the tensile strength was measured on all specimens that reached matrix is chemically unstable and forms mullite at 1300C.31,32 run-out during fatigue and creep testing Below 1200C, the diphasic Al, osio, mixture is expected to be (D) Interlaminar Strength: A compressive load was applie kinetically stable with respect to mullite formation. 1 32 No evi- to a notched specimen of uniform width using ASTM standard test dence of mullite formation in the matrix has been found using practice D3846("Standard Test Method for In-Plane Shear teM or XRd Book of standards. Vol 08.02 ASTM International. West Con- shohocken, PA). The specimens failed in shear between two (2) Microstructure and Residual Stress entrally located notches machined halfway through the spec No interphase was present between the matrix and fiber. The thickness at a standard distance apart on opposing faces. Tests same porous SiO, that bonded matrix Al,O, grains together also were done in stroke control with a rate of I mm/min at 23. 538 bonded those grains to the Nextel 610 fibers(Fig. 2). The Nextel 610 fiber 0.11±0.03 982°and1037° In air distributed uniformly through the fiber, without preference for (E) Cyclic Tension: Cyclic tension (fatigue) tests were intragranular or intergranular location. The average pore diameter rformed on dogbone test specimens at room temperature and was 9.8+ 4.2 nm. Fiber porosity did not coarsen after 3000 h at 1000%C. The tests were conducted in load control with a load ratio 982°C(1800°F) of 0.05(R=min/oma ). Room-temperature tests were cycled at 1 Hz for the first 100 000 cycles and then at 5 Hz for an additional 900 000 cycles, or until failure. The 1000C fatigue tests were done at a frequency of I Hz and were allowed to run for 100 000 cycles. The 1000oC cycle count value was chosen to roughly duplicate the number of loadings expected in aerospace applica- ns at that temperature. Fatigue run-out limits were defined to be the stress level at or slightly above the highest run-out stress for the g-N relationships. Changes in hysteresis energy density and elastic modulus versus fatigue cycle were calculated for each fatigue test to assess the extent of damage that occurred to the composite Creep Rupture: Creep rupture tests on dogbone test ,pecimens were conducted under load control at 75,100, 125, and MPa at 1000C. The 1100%C tests were performed at 50 and 75 MPa. The run-out condition was defined as 100 h for both test temperatures. The specimens were ramped at 10 MPa/s to the est stress. Data were recorded during loading up to the test load and from the time the specimen reached the test stress so that total 腾 strain and creep strain could be calculated I. Results and discussion N6As装120m (1) General Microstructure Fig. 1. Optical micrograph of Nextel 610/AS showing Fiber distribution in the matrix was not uniform The matrix was taken 45 to the warp direction of the fibers. Micrograph shows matrix-rich a homogenous mixture of a-Al,O, grains cemented together by regions between plies, matrix cracks, and fiber distribution
(B) Mechanical Test Apparatus: A horizontal servohydraulic machine with rigid hydraulic clamping grips and quartz-lamp heating was used for the tension, in-plane shear, creep rupture, and fatigue tests. Test control, data acquisition, and interactive data analysis was done using the MATE program24 on an IBM-compatible personal computer (PC) linked to the test frame by an analog-to-digital board. Temperature was measured using five S-type thermocouples bonded to each specimen with an alumina-based ceramic adhesive. A detailed description of the test equipment was given elsewhere.25 The interlaminar strength and in-plane shear tests were done using a standard vertical servohydraulic test machine with a box furnace that used igniter elements. Monotonic tests consisted of testing two to three test specimens per condition. Fatigue and creep rupture testing consisted of testing only one test specimen at each stress level investigated. It was recognized by the authors that this was an extremely limited set of data on which to make scientific observations. However, the objective was to explore the boundaries of mechanical behavior for oxide/oxide CMCs and to determine if they warranted a more rigorous investigation for use in aerospace applications. (C) Monotonic Loading: All tension and compression tests were performed using stroke control with a 0.05 mm/s displacement rate. Tension tests were performed using dogbone test specimens, whereas the compression tests used straight-sided test specimens. Tension tests were also done at 1000° and 1100°C in air with an 1.5 cm hot zone. In high-temperature tests, each specimen was ramped to the test temperature in 15 min and then equilibrated for 20 min; the stress then was ramped up until the specimen failed. For in-plane shear measurement, tension tests were performed using test specimens with 45° fiber orientations. Residual room-temperature tensile strength was measured on all specimens that reached run-out during fatigue and creep testing. (D) Interlaminar Strength: A compressive load was applied to a notched specimen of uniform width using ASTM standard test practice D3846 (“Standard Test Method for In-Plane Shear Strength of Reinforced Plastics,” Designation No. D3846, ASTM Book of Standards, Vol. 08.02, ASTM International, West Conshohocken, PA). The specimens failed in shear between two centrally located notches machined halfway through the specimen thickness at a standard distance apart on opposing faces. Tests were done in stroke control with a rate of 1 mm/min at 23°, 538°, 982°, and 1037°C in air. (E) Cyclic Tension: Cyclic tension (fatigue) tests were performed on dogbone test specimens at room temperature and 1000°C. The tests were conducted in load control with a load ratio of 0.05 (R min/max). Room-temperature tests were cycled at 1 Hz for the first 100 000 cycles and then at 5 Hz for an additional 900 000 cycles, or until failure. The 1000°C fatigue tests were done at a frequency of 1 Hz and were allowed to run for 100 000 cycles. The 1000°C cycle count value was chosen to roughly duplicate the number of loadings expected in aerospace applications at that temperature. Fatigue run-out limits were defined to be the stress level at or slightly above the highest run-out stress for the –N relationships. Changes in hysteresis energy density and elastic modulus versus fatigue cycle were calculated for each fatigue test to assess the extent of damage that occurred to the composite. (F) Creep Rupture: Creep rupture tests on dogbone test specimens were conducted under load control at 75, 100, 125, and 135 MPa at 1000°C. The 1100°C tests were performed at 50 and 75 MPa. The run-out condition was defined as 100 h for both test temperatures. The specimens were ramped at 10 MPa/s to the test stress. Data were recorded during loading up to the test load and from the time the specimen reached the test stress so that total strain and creep strain could be calculated. III. Results and Discussion (1) General Microstructure Fiber distribution in the matrix was not uniform. The matrix was a homogenous mixture of -Al2O3 grains cemented together by porous SiO2. A micrograph taken 45° to the warp and fill fibers is shown in Fig. 1. Parallel arrays of cracks in the matrix were perpendicular to the cloth layers. The crack spacing was wider in the matrix-rich regions and smaller within the fiber tows. Presumably, these cracks formed by matrix shrinkage during sintering that was constrained by the cloth layers, similar to that observed for constrained sintering of films26 and around inclusions.27–29 These cracks formed under relatively low stresses that could not exceed twice the sintering stress.30 The fiber–matrix interfaces and the fibers themselves were not cracked. The average skeletal densities of seven specimens were 3.62 g/cm3 using the Archimedes method and 3.65 g/cm3 using the pycnometer, with a standard deviation of 0.06. An average bulk density of 2.90 g/cm3 was calculated from the immersion measurements. The average fiber volume fraction was 30.7% 2.29%. The SiO2 and Al2O3 volume fractions of the matrix were not measured, but the expected proportions following processing were 87 wt% -Al2O3 and 13 wt% amorphous SiO2 (2.2 g/cm3 ). Therefore, the interconnected porosity in the entire composite was 24%, or 35% in the matrix alone. Al2O3 occupied 51 vol% of the matrix and SiO2 14%. Because this matrix porosity was concentrated in the SiO2, initial SiO2 porosity was 71%. However, sintering shrinkage cracks accounted for some of this porosity; the remainder was finely distributed in the SiO2 that cemented the Al2O3 grains together. The high specific surface area measured by BET nitrogen adsorption of 25 to 35 m2 /g (average of 31.25 3.98 m2 /g) was consistent with a large amount of fine interconnected porosity. In contrast to N610/AS, dense glass-ceramic composites of Nicalon/MAS were found to have a specific surface area of 0.3 m2 /g. The mixture of -Al2O3 and amorphous SiO2 that forms the matrix is chemically unstable and forms mullite at 1300°C.31,32 Below 1200°C, the diphasic Al2O3–SiO2 mixture is expected to be kinetically stable with respect to mullite formation.31,32 No evidence of mullite formation in the matrix has been found using TEM or XRD. (2) Microstructure and Residual Stress No interphase was present between the matrix and fiber. The same porous SiO2 that bonded matrix Al2O3 grains together also bonded those grains to the Nextel 610 fibers (Fig. 2). The Nextel 610 fiber grain size was 0.11 0.03 m. Fine porosity was distributed uniformly through the fiber, without preference for intragranular or intergranular location. The average pore diameter was 9.8 4.2 nm. Fiber porosity did not coarsen after 3000 h at 982°C (1800°F). Fig. 1. Optical micrograph of Nextel 610/AS showing a cross section taken 45° to the warp direction of the fibers. Micrograph shows matrix-rich regions between plies, matrix cracks, and fiber distribution. 982 Journal of the American Ceramic Society—Zawada et al. Vol. 86, No. 6
June 2003 Characteriation and High-Temperature Mechanical Behavior of an Oxide/Oxide Composite No Heat- Treatment to Heaf-Treatment Fiber Fig. 2. High-resolution TEM micrograph showing the complete absence of an interphase between the fiber and matrix in Nextel 610/AS 3000 hours a982℃ (b) In as-fabricated specimens, the matrix AL, O, grains were 0 19+0.08 um in diameter, with an average aspect ratio of 1.54 The Sio, between the Al2O, grains was amorphous with a large volume fraction of mostly interconnected porosity(Fig. 3).The SiO, pore volume fraction was so high that it was unclear whether the radii relevant for coarsening analysis were those of the pores those of the SiO,. The pore radii had a lognormal distribution with an inverse log average of 0.58 0.18 log(nm)(Fig 4(a)). In specimens heat-treated for 3000 h at 982%C(1800%F), the Al,O3 matrix grains averaged 0. 17+ 0.07 um in diameter with an aspect ratio of 1. 43. Statistically, there was no change in AlO, grain size or shape with heat treatment at 982.C. However, the average po radi in the Sioz increased by a large amount (Fig 4(b). The SiOz nm ore radii again had a lognormal distribution, but with an inverse average of 1.04 0.31 log(nm). In some cases the SiO2 formed a dense coating around Al,O3 grains(Fig 4(b)) Fig. 4.(a) TEM micrograph of as-processed Nextel 610/AS matrix 时m in the matrix SiO2.(b)TEM micrograph of Nextel 610/AS ning of the porosity in the matrix SiO2 after No Heat-Treatine rate can be an appro imate SiO, viscosity. The pores are interconnected and the pore radii distributions are lognormal rather than large-sized skewed ormal distributions typical for steady-state coarsening. This that a"cylinder"model for the porosity">, is appropri e tightly packed Al2O3 grains that surround the Sio, and the constrain the porous SiO2 not allow densification. >Therefore coarsening occurs instead of A pore-coarsening calculation using the as-fabricated and 3000 h/982C coarsening data has been done by scherer. 3> The calcu lation assumes a SiO, surface energy of 0.28 J/m2.36 A Sio viscosity of 1.1 x 10 Pas is calculated, which is about the viscosity expected for SiO, with 0.04%0.12% hydrox nd from the n "16 MPa. Because, in this case, sintering is constrained, the sintering stress is a real tensile stress in the porous SiO, supported by a small compressive stress at the Al,OAl,O, grain contacts in the matrix, The sintering stress can be present only at high temperatures when mass transport processes redistribute SiO, over the 0.1 um between Al,O, grains A rough estimation of the matrix stress state now can be made The high CTE difference(8X 10-6oC-)and modulus difference between Al,O3 and porous SiO, should cause 200 MPa residual hydrostatic compression in the porous SiO, at room temperature. 3. High-resolution TEM micrograph on Nextel 610/AS showing the balanced by a tensile stress at AL,OxAL, O, grain contacts al structure of the matrix. Matrix is made up of single-crystal Al,O Thermal contraction of the fiber plies superimposes compressive es suspended in SiO2 that contains very fine porosity plane stress on this matrix residual stress. Residual stresses vanish
In as-fabricated specimens, the matrix Al2O3 grains were 0.19 0.08 m in diameter, with an average aspect ratio of 1.54. The SiO2 between the Al2O3 grains was amorphous with a large volume fraction of mostly interconnected porosity (Fig. 3). The SiO2 pore volume fraction was so high that it was unclear whether the radii relevant for coarsening analysis were those of the pores or those of the SiO2. The pore radii had a lognormal distribution with an inverse log average of 0.58 0.18 log(nm) (Fig. 4(a)). In specimens heat-treated for 3000 h at 982°C (1800°F), the Al2O3 matrix grains averaged 0.17 0.07 m in diameter with an aspect ratio of 1.43. Statistically, there was no change in Al2O3 grain size or shape with heat treatment at 982°C. However, the average pore radii in the SiO2 increased by a large amount (Fig. 4(b)). The SiO2 pore radii again had a lognormal distribution, but with an inverse log average of 1.04 0.31 log(nm). In some cases the SiO2 formed a dense coating around Al2O3 grains (Fig. 4(b)). The pore-coarsening rate can be analyzed to obtain an approximate SiO2 viscosity. The pores are interconnected and the pore radii distributions are lognormal rather than large-sized skewed normal distributions typical for steady-state coarsening. This suggests that a “cylinder” model for the porosity33,34 is appropriate. The tightly packed Al2O3 grains that surround the SiO2 and the fibers that surround the matrix constrain the porous SiO2 and do not allow densification.35 Therefore, coarsening occurs instead of sintering. A pore-coarsening calculation using the as-fabricated and 3000 h/982°C coarsening data has been done by Scherer.35 The calculation assumes a SiO2 surface energy of 0.28 J/m2 . 36 A SiO2 viscosity of 1.1 1014 Pas is calculated, which is about the viscosity expected for SiO2 with 0.04%–0.12% hydroxyl groups.37 The “sintering stress”38 found from the same calculation is 16 MPa. Because, in this case, sintering is constrained, the sintering stress is a real tensile stress in the porous SiO2 supported by a small compressive stress at the Al2O3–Al2O3 grain contacts in the matrix. The sintering stress can be present only at high temperatures when mass transport processes redistribute SiO2 over the 0.1 m spaces between Al2O3 grains. A rough estimation of the matrix stress state now can be made. The high CTE difference (8 10–6 °C–1 ) and modulus difference between Al2O3 and porous SiO2 should cause 200 MPa residual hydrostatic compression in the porous SiO2 at room temperature, balanced by a tensile stress at Al2O3–Al2O3 grain contacts. Thermal contraction of the fiber plies superimposes compressive plane stress on this matrix residual stress. Residual stresses vanish Fig. 2. High-resolution TEM micrograph showing the complete absence of an interphase between the fiber and matrix in Nextel 610/AS. Fig. 3. High-resolution TEM micrograph on Nextel 610/AS showing the typical structure of the matrix. Matrix is made up of single-crystal Al2O3 particles suspended in SiO2 that contains very fine porosity. Fig. 4. (a) TEM micrograph of as-processed Nextel 610/AS matrix showing the very fine porosity in the matrix SiO2. (b) TEM micrograph of Nextel 610/AS showing coarsening of the porosity in the matrix SiO2 after heat treatment at 982°C for 3000 h. June 2003 Characterization and High-Temperature Mechanical Behavior of an Oxide/Oxide Composite 983
984 Journal of the American Ceramic Sociery-Zawada et al. Vol 86. No 6 as the composite is brought to the processing temperature (1000C), where, with time, a negligible tensile sintering stress in SiO, may develop Load Rate = 0.05 mm/s OUTs"205 MPa ( Monotonic Loading Tests A) Tension: The results of the monotonic tension tests are 三150 shown in Figs. 5(aHc). The stress-strain response is nearly linear-elastic to failure for all specimens and all temperature investigated. Such linear behavior suggests that there is little 10 additional matrix cracking during loading and that fiber-matrix debonding is insignificant. In tension, fiber fracture appears to be the dominant damage mode and is typical fiber-dominated com osite behavior. The average room-temperature ultimate tensile trength and strain to failure are 205 MPa and 0.3%, respectively Fig. 5(a)). The average room-temperature elastic modulus is -70 GPa. Such a value is significantly lower than the value of 200 GPa measured for traditional SiC-fiber CMCs of the same fiber volume Strain (% fraction. High-temperature tensile strength decreases only 15% from room temperature, and the modulus is relatively unchanged The short-term tensile behavior of this composite does not change significantly for temperatures up to 1 100C The fracture surfaces were examined using low-magnification Load Rate = 0.05 mm/s optical microscopy. The room-temperature specimens had irregu- lar fracture paths and"fiber bundle pullout"with matrix material aurs≈173MPa E150 they were significantly less common in the high-temperature tests (Fig. 6(b). Fiber bundle pullout was not caused by the fracture mechanisms that caused single-fiber pullout in CMCs with weak 100 fiber-matrix interfaces. The apparent fiber bundle pullout was cpL85 MPa simply due to weak matrix material falling apart during failure observed in tension tests of other CMCs with the same fiber 40 E。=80GPa n One very important observation was that the room-temperature become fracture surfaces were substantially more jagged. The crack path deviated substantially across the width as well as along the length of the specimen, and the fracture surface extended -10 mm along Strain (% the length of the gauge section. The fracture paths stepped across the specimen in the thickness direction, and some of the bundle lengths were several millimeters in length. In contrast, the elevated-temperature fracture surfaces were significantly more lanar, with only very short bundle lengths evident on the fracture Load Rate 0.05 mm/s surfaces. In high-temperature tests, the fracture surface was con- 71 MPa fined to only 3-4 mm along the specimen gauge length. The distinct change in fracture path morphology suggested a funda- mental change in damage mechanisms with temperature. It was uggested that this change, in part, may have been due to the sio2 in the matrix and how residual stress changed with increased temperature, as discussed earlier. The stress versus strain behavior of this N610/AS composite is very similar to composites made by other manufacturers. A Nextel 610 composite with a matrix of 80% mullite and 20% Al,O3 demonstrates tensile behavior similar to what has been measured E.≈75GPa ecimen #2 in this investigation. Tensile strengths were found to be 200 MPa, and the stress versus strain traces were essentially linear to failure. The room-temperature tensile properties of several oxide/ oxide CMCs were measured by one of the authors in a separate Strain(%) investigation. Each of the oxide/oxide systems investigated had similar stress versus strain traces that were essentially linear to Fig. 5.(a)Tensile behavior for three test specimens of failure. Differences in modulus and ultimate strength were primar- xtel 610/AS tension tested at 23C.(b) Tensile stress ily attributed to the volume fraction of fibers used and type of avior for two test specimens of &HS W Nextel 610/AS tension fiber. A review of the mechanical properti porous-matrIx 0oC.(c)Tensile stress versus strain behavior of &HS W Nexte ceramic composites was given by Zok and Levi. One important statement in the review was that because of the inherent nature of matrix. much of the stress-strain behavior was controlled by the fiber and that ultimate strength was a Fig. 7 as shear stress versus shear strain. The stress-strain respons function of how much the fibers were damaged ear at -50% of ultimate B) In-Plane Shear and Interlaminar Strength: Stress strength of 27 GPa was observed for three tests. On the specimen strain response in room-temperature in-plane shear testing or faces. there was a dominant shear band at a 45 angle across the specimens with +45 fiber orientations was very different from gauge section that was 10 mm wide(-7-10 fiber tows). At the specimens with 0/90 fiber orientations. The results are plotted in edge of the failed specimens, the fiber tows along the shear band
as the composite is brought to the processing temperature (1000°C), where, with time, a negligible tensile sintering stress in SiO2 may develop. (3) Monotonic Loading Tests (A) Tension: The results of the monotonic tension tests are shown in Figs. 5(a)–(c). The stress–strain response is nearly linear-elastic to failure for all specimens and all temperatures investigated. Such linear behavior suggests that there is little additional matrix cracking during loading and that fiber–matrix debonding is insignificant. In tension, fiber fracture appears to be the dominant damage mode and is typical fiber-dominated composite behavior. The average room-temperature ultimate tensile strength and strain to failure are 205 MPa and 0.3%, respectively (Fig. 5(a)). The average room-temperature elastic modulus is 70 GPa. Such a value is significantly lower than the value of 200 GPa measured for traditional SiC-fiber CMCs of the same fiber volume fraction. High-temperature tensile strength decreases only 15% from room temperature, and the modulus is relatively unchanged. The short-term tensile behavior of this composite does not change significantly for temperatures up to 1100°C. The fracture surfaces were examined using low-magnification optical microscopy. The room-temperature specimens had irregular fracture paths and “fiber bundle pullout” with matrix material remaining on the fiber surfaces. These were common fracture features in all the room-temperature tension tests (Fig. 6(a)), but they were significantly less common in the high-temperature tests (Fig. 6(b)). Fiber bundle pullout was not caused by the fracture mechanisms that caused single-fiber pullout in CMCs with weak fiber–matrix interfaces.39 The apparent fiber bundle pullout was simply due to weak matrix material falling apart during failure; fracture surfaces did not mate. Fiber failure was intergranular, as observed in tension tests of other CMCs with the same fiber.40 One very important observation was that the room-temperature fracture surfaces were substantially more jagged. The crack path deviated substantially across the width as well as along the length of the specimen, and the fracture surface extended 10 mm along the length of the gauge section. The fracture paths stepped across the specimen in the thickness direction, and some of the bundle lengths were several millimeters in length. In contrast, the elevated-temperature fracture surfaces were significantly more planar, with only very short bundle lengths evident on the fracture surfaces. In high-temperature tests, the fracture surface was confined to only 3–4 mm along the specimen gauge length. The distinct change in fracture path morphology suggested a fundamental change in damage mechanisms with temperature. It was suggested that this change, in part, may have been due to the SiO2 in the matrix and how residual stress changed with increased temperature, as discussed earlier. The stress versus strain behavior of this N610/AS composite is very similar to composites made by other manufacturers. A Nextel 610 composite with a matrix of 80% mullite and 20% Al2O3 demonstrates tensile behavior similar to what has been measured in this investigation.41 Tensile strengths were found to be 200 MPa, and the stress versus strain traces were essentially linear to failure. The room-temperature tensile properties of several oxide/ oxide CMCs were measured by one of the authors in a separate investigation. Each of the oxide/oxide systems investigated had similar stress versus strain traces that were essentially linear to failure. Differences in modulus and ultimate strength were primarily attributed to the volume fraction of fibers used and type of fiber. A review of the mechanical properties of porous-matrix ceramic composites was given by Zok and Levi.22 One important statement in the review was that, because of the inherent nature of a porous low-energy matrix, much of the stress–strain behavior was controlled by the fiber and that ultimate strength was a function of how much the fibers were damaged. (B) In-Plane Shear and Interlaminar Strength: Stress– strain response in room-temperature in-plane shear testing on specimens with 45° fiber orientations was very different from specimens with 0°/90° fiber orientations. The results are plotted in Fig. 7 as shear stress versus shear strain. The stress–strain response became nonlinear at 50% of ultimate stress. An average shear strength of 27 GPa was observed for three tests. On the specimen faces, there was a dominant shear band at a 45° angle across the gauge section that was 10 mm wide (7–10 fiber tows). At the edge of the failed specimens, the fiber tows along the shear band Fig. 5. (a) Tensile stress versus strain behavior for three test specimens of 8HSW Nextel 610/AS tension tested at 23°C. (b) Tensile stress versus strain behavior for two test specimens of 8HSW Nextel 610/AS tension tested at 1000°C. (c) Tensile stress versus strain behavior of 8HSW Nextel 610/AS tested at 1100°C. 984 Journal of the American Ceramic Society—Zawada et al. Vol. 86, No. 6
Characteriation and High-Temperature Mechanical Behavior of an Oxide/Oxide Composite 985 Load Rate 94311 三n0苏 30 20 TpL 17MPa 10 G k 16 GPa Fig. 7. Shear stress versus shear strain at room t as tested using a±45° fiber or 10 mm increased slightly with temperature(Fig 8), while the ar shear strength increased significantly the test temp ncreased( Fig. 9). The low in-plane and interlaminar failure fi ngth should have made this composite less sensitive to failure from center or edge notches -Increased shea strengths with increased temperature again suggested that th viscosity of the matrix SiO, and residual stress state may have affected the observed behavior (4 Cyclic Tension Results of the room-temperature and 1000C fatigue tests are shown in Figs. 10(a)and(b), respectively. The room-temperature fatigue limit is 170 MPa, which is -85% of the average ature tensile strength. The 1000C fatigue limit is 150 MPa, which is -85% of the average 1000.C tensile strength. The room-temperature fatigue performance is similar to other CMCs, many of which exhibit fatigue limits within 5%20% of the 10 mn average tensile strength. However, the fatigue behavior at elevated temperature for this CMC is unlike that observed for any other CMC. CMCs with an interphase consisting of carbon or BN typically exhibit run-out at stress levels of only 75-120 MPa and these run-out stress levels are always closely associated with the proportional limit and development of matrix crack ing s Once the matrix is cracked, there is rapid on o the oxygen into the composite, resulting in oxidat Fig. 6.(a) Fracture surface of Nextel 610/AS tension tested at 23.C Fracture surface shows extensive damage. a damage zone -15 length, and macrofeatures associated with the fiber tow bundles. Fracture surface of Nextel 610/as tension tested at 1000oC. Fracture surface is very flat and is confined to a very narrow zone along the length ±45 Tension Test of the specimen were pulled I mm into the composite, while the fiber tows to the shear band remained in of the There was extensive matrix matrix allowed the tows to rotate as they withdrew. This may have promoted nonlinear stress-strain behavior beyond the onset of strain localization. The shear strengths were very low compared with the 0/90%orienta- tion and several times lower than those of Nicalon/SiC Cmcs (100 MPa) measured using the losipescu test fixture. Very similar behavior was observed for a N610/mullite Al,O3 CMC tested using the +45 fiber orientation coupon loaded in tension. 4 Maximum stress levels were measured to be 63 MPa, whereas Temperature (C) this investigation measured an average failure stress of 54 MPa. Photographs of a fractured *450 tensile specimen were nearly Fig. 8. Interlaminar shear strength versus temperature for 8HSW Ne identical to the fracture features produced in this investigation. were generated using a double- notched com
were pulled 1 mm into the composite, while the fiber tows 90° to the shear band remained in place. There was extensive matrix damage. The fragmentation of the matrix allowed the tows to rotate as they withdrew. This may have promoted nonlinear stress–strain behavior beyond the onset of strain localization. The shear strengths were very low compared with the 0°/90° orientation and several times lower than those of Nicalon/SiC CMCs (100 MPa) measured using the Iosipescu test fixture.42 Very similar behavior was observed for a N610/mullite Al2O3 CMC tested using the 45° fiber orientation coupon loaded in tension.41 Maximum stress levels were measured to be 63 MPa, whereas this investigation measured an average failure stress of 54 MPa. Photographs of a fractured 45° tensile specimen were nearly identical to the fracture features produced in this investigation. The in-plane strength increased slightly with temperature (Fig. 8), while the interlaminar shear strength increased significantly as the test temperature increased (Fig. 9). The low in-plane and interlaminar shear strength should have made this composite less sensitive to failure from center or edge notches.43 Increased shear strengths with increased temperature again suggested that the viscosity of the matrix SiO2 and residual stress state may have affected the observed behavior. (4) Cyclic Tension Results of the room-temperature and 1000°C fatigue tests are shown in Figs. 10(a) and (b), respectively. The room-temperature fatigue limit is 170 MPa, which is 85% of the average room-temperature tensile strength. The 1000°C fatigue limit is 150 MPa, which is 85% of the average 1000°C tensile strength. The room-temperature fatigue performance is similar to other CMCs, many of which exhibit fatigue limits within 5%–20% of the average tensile strength. However, the fatigue behavior at elevated temperature for this CMC is unlike that observed for any other CMC. CMCs with an interphase consisting of carbon or BN typically exhibit run-out at stress levels of only 75–120 MPa, and these run-out stress levels are always closely associated with the proportional limit and development of matrix cracking.18 Once the matrix is cracked, there is rapid ingress of oxygen into the composite, resulting in oxidation of the Fig. 6. (a) Fracture surface of Nextel 610/AS tension tested at 23°C. Fracture surface shows extensive damage, a damage zone 15 mm in length, and macrofeatures associated with the fiber tow bundles. (b) Fracture surface of Nextel 610/AS tension tested at 1000°C. Fracture surface is very flat and is confined to a very narrow zone along the length of the specimen. Fig. 7. Shear stress versus shear strain at room temperature for 8HSW Nextel 610/AS tested using a 45° fiber orientation in the tensile specimen. Fig. 8. Interlaminar shear strength versus temperature for 8HSW Nextel 610/AS. Data were generated using a double-notched compression specimen. June 2003 Characterization and High-Temperature Mechanical Behavior of an Oxide/Oxide Composite 985
986 Journal of the American Ceramic Sociery-Zawada et al. Vol 86. No 6 interphase between the fiber and matrix and loss in strength. In causes changes in elastic modulus and stress-strain hysteresis, %R4 as contrast, the high-temperature fatigue life of this oxide/o CMC does not substantially decrease with applied stress Continuous degradation of fatigue life in CMCs is usua observed to depend on applied stress, stress ratio, frequency, and test environments. Matrix damage, fiber debonding, and fibe sliding are usually observed during cyclic loading. Fatigue damag 10 However, for N610/AS, the longitudinal modulus does not appear to change significantly. The normalized modulus values for the 1000.C fatigue tests are shown in Fig. 11. On the first few cycles, there is a 5%0-10% decrease in modulus. after which the stiffness remains constant out to 1000 cycles. From this cycle count on out to 100 000 cycles, there appears to be only a slight decrease in stiffness with continued cycling. This suggests that there is 1000 little progressive damage with continued cycling, and most likely this is limited to slight additional matrix cracking. The hysteresis Temperature(°c) energy density behavior mirrors the modulus behavior. However, the values for hysteretic energy density are extremely small, as Nextel 610/AS. Data were generated using a shown in Fig. 12, with an average hEd value of only 3-5 kJ/m3 nina shear strength versus temperature at room Traditional composites with interphases and classical fiber debonding typically generate HED values of 280 KJ/m when fatigued above the proportional limit. For stress levels between 100 and 150 MPa, there is little difference in the hed behavior for N610/AS. Such small values of HED indicate that there is very little hysteresis and suggests that little actual fatigue damage Another way to qualify fatigue damage is to monitor strain. 13 plots maximum and minimum strain versus cycles for 250 the fatigue test conducted at 150 MPa and 1000C. There is very ension Test little evidence of strain accumulation for the 100 000 cycles tested at23°c 200严 In addition, the difference between maximum and minimum strain for a given cycle does not change. N610/AS appears to neither Fatigue Limit cyclic strain soften nor harden. This further substantiates that little 2 170 MPa fatigue damage has occurred for this test. 150 After the fatigue tests were completed, the specimens were cooled to room temperature and tension tested to measure the 100 etained strength. Measuring retained strength was important for determining the damage state of the specimen, especially if emperature tensile strengths of those simens that reached run-out during fatigue testing are shown in Table I. In most CMcs (a) d to hi 10°1011021031 run-out. However. for this composite. the tensile strength was not decreased. The specimen fatigued at 1000.C for 100 000 cycles at 150 MPa had no loss of tensile strength, and the stress-strain curve 250 110 200 100 日 03o9N=Eoz 70 =Rf 10101102103104 Cycles(N) Cycles(n) Fig 10. (a) Fatigue plot for 8HSW Nextel 610/AS showing stress ve cles to failure for fatigue tests at room temperature. (b) Fatigue plot for Fig. 11. Plot of normalized modulus versus fatigue cycles for Nextel 8HSW Nextel 610/As showing stress versus cycles to failure for fatigue 610/AS tested at 1000%C and four stress levels. At each stress level there tests at I000°C is a slight decrease in modulus with increased cycle count
interphase between the fiber and matrix and loss in strength. In contrast, the high-temperature fatigue life of this oxide/oxide CMC does not substantially decrease with applied stress. Continuous degradation of fatigue life in CMCs is usually observed to depend on applied stress, stress ratio, frequency, and test environments. Matrix damage, fiber debonding, and fiber sliding are usually observed during cyclic loading. Fatigue damage causes changes in elastic modulus and stress–strain hysteresis.43,44 However, for N610/AS, the longitudinal modulus does not appear to change significantly. The normalized modulus values for the 1000°C fatigue tests are shown in Fig. 11. On the first few cycles, there is a 5%–10% decrease in modulus, after which the stiffness remains constant out to 1000 cycles. From this cycle count on out to 100 000 cycles, there appears to be only a slight decrease in stiffness with continued cycling. This suggests that there is very little progressive damage with continued cycling, and most likely this is limited to slight additional matrix cracking. The hysteresis energy density behavior mirrors the modulus behavior. However, the values for hysteretic energy density are extremely small, as shown in Fig. 12, with an average HED value of only 3–5 kJ/m3 . Traditional composites with interphases and classical fiber debonding typically generate HED values of 80 KJ/m3 when fatigued above the proportional limit. For stress levels between 100 and 150 MPa, there is little difference in the HED behavior for N610/AS. Such small values of HED indicate that there is very little hysteresis and suggests that little actual fatigue damage should develop. Another way to qualify fatigue damage is to monitor strain. Figure 13 plots maximum and minimum strain versus cycles for the fatigue test conducted at 150 MPa and 1000°C. There is very little evidence of strain accumulation for the 100 000 cycles tested. In addition, the difference between maximum and minimum strain for a given cycle does not change. N610/AS appears to neither cyclic strain soften nor harden. This further substantiates that little fatigue damage has occurred for this test. After the fatigue tests were completed, the specimens were cooled to room temperature and tension tested to measure the retained strength. Measuring retained strength was important for determining the damage state of the specimen, especially if oxidation had occurred during testing. The retained roomtemperature tensile strengths of those specimens that reached run-out during fatigue testing are shown in Table I. In most CMCs exposed to high-temperature fatigue testing, degradation of interface properties decreases tensile strength in specimens that reach run-out. However, for this composite, the tensile strength was not decreased. The specimen fatigued at 1000°C for 100 000 cycles at 150 MPa had no loss of tensile strength, and the stress–strain curve Fig. 9. Tensile interlaminar shear strength versus temperature at room temperature for 8HSW Nextel 610/AS. Data were generated using a double-notch compression specimen. Fig. 10. (a) Fatigue plot for 8HSW Nextel 610/AS showing stress versus cycles to failure for fatigue tests at room temperature. (b) Fatigue plot for 8HSW Nextel 610/AS showing stress versus cycles to failure for fatigue tests at 1000°C. Fig. 11. Plot of normalized modulus versus fatigue cycles for Nextel 610/AS tested at 1000°C and four stress levels. At each stress level there is a slight decrease in modulus with increased cycle count. 986 Journal of the American Ceramic Society—Zawada et al. Vol. 86, No. 6
June 2003 Characteriation and High-Temperature Mechanical Behavior of an Oxide/Oxide Composite Table L. Measured Residual Strength of Samples that Reached Run-Out during Fatigue and en #4, 150MPa Test stress level Residual strength k一 P All residual strength tests were conducted at room temperature. (5) Creep Rupture The results of the creep strain versus time for the 1000.C tests Cycles (N are shown in Fig. 14(a), while the 1.C results are shown in Fig 14(b). The primary creep regime is extremely short and transitions ersus fatigue cycles for Nextel apidly into secondary creep. Secondary creep is essentially linear our stress levels. At each stress level there is a ing absorbed by the specimen, and until failure. No tertiary creep is observed. In all cases, as the after -1000 applied stress level increases, the strain to failure decreases and the strain rate increases. Figure 14(a) shows that, for stress levels of 2125 MPa, the strain to failure is roughly equal to the strain during the tension tests and suggests the mechanism controlling as nearly identical to the tensile stress-strain curves shown in damage is rate sensitive. A plot of applied creep rupture stress Fig. 5(a). The only change was that the stress-strain trace was level versus time to failure for both temperatures is shown in Fig completely linear. Fatigue cycling appeared to have worked out all 15. The run-out stress is 75 MPa for creep at 1000.C.However, th the nonlinear behavior by saturating the specimen with matrix creep resistance at 1100C is so poor that run-out is not even cracks. The residual strength measurement suggested that no observed for stresses as low as 50 MPa. The full scale for time in fatigue damage occurred to the fibers during these fatigue tests and suggested the small decrease in stiffness was most likely associ- ited with some slight additional progressive matrix cracking The excellent fatigue resistance shown by the N610/AS CMC studied in this investigation is most likely attributed to the porous T=1000°c matrix. There is little high-temperature durability data on porous- 75 MPa matrix CMCs, but the excellent fatigue behavior of the porous 135 MPa matrix CMC investigated in this study does extend to another 125 MPa porous-matrix oxide/oxide CMC system. Steel" has investigated the fatigue behavior of &HSW N720/Al,O3 and has found the 100 MPa room-temperature and 1200C fatigue limits to be within 95% of a the measured average tensile strengths at those temperatures. Run-out test specimens have been tested for retained strength, and their strengths are also found to be noticeably higher than the 25 two different orous-matrix oxide/oxide systems that there is no fatigue-induced degradation in fiber strength during fatigue and that the retained 100000 trengths are greater than the as-received strength Time (s) T=1100c d20 5,000 1010210310410510 Time(s) Cycles Fig. 14. (a) Creep time for Nextel 6 tested at
was nearly identical to the tensile stress–strain curves shown in Fig. 5(a). The only change was that the stress–strain trace was completely linear. Fatigue cycling appeared to have worked out all the nonlinear behavior by saturating the specimen with matrix cracks. The residual strength measurement suggested that no fatigue damage occurred to the fibers during these fatigue tests and suggested the small decrease in stiffness was most likely associated with some slight additional progressive matrix cracking. The excellent fatigue resistance shown by the N610/AS CMC studied in this investigation is most likely attributed to the porous matrix. There is little high-temperature durability data on porousmatrix CMCs, but the excellent fatigue behavior of the porousmatrix CMC investigated in this study does extend to another porous-matrix oxide/oxide CMC system. Steel45 has investigated the fatigue behavior of 8HSW N720/Al2O3 and has found the room-temperature and 1200°C fatigue limits to be within 95% of the measured average tensile strengths at those temperatures. Run-out test specimens have been tested for retained strength, and their strengths are also found to be noticeably higher than the as-received strength. There is now evidence in two different porous-matrix oxide/oxide systems that there is no fatigue-induced degradation in fiber strength during fatigue and that the retained strengths are greater than the as-received strengths. (5) Creep Rupture The results of the creep strain versus time for the 1000°C tests are shown in Fig. 14(a), while the 1100°C results are shown in Fig. 14(b). The primary creep regime is extremely short and transitions rapidly into secondary creep. Secondary creep is essentially linear until failure. No tertiary creep is observed. In all cases, as the applied stress level increases, the strain to failure decreases and the strain rate increases. Figure 14(a) shows that, for stress levels of 125 MPa, the strain to failure is roughly equal to the strain during the tension tests and suggests the mechanism controlling damage is rate sensitive. A plot of applied creep rupture stress level versus time to failure for both temperatures is shown in Fig. 15. The run-out stress is 75 MPa for creep at 1000°C. However, the creep resistance at 1100°C is so poor that run-out is not even observed for stresses as low as 50 MPa. The full scale for time in Fig. 12. Hysteretic energy density versus fatigue cycles for Nextel 610/AS fatigue tested at 1000°C and four stress levels. At each stress level there is a rapid decrease in energy being absorbed by the specimen, and, after 1000 cycles, the value remains constant. Fig. 13. Maximum and minimum percent strain measured during a fatigue test of Nextel 610/AS tested at 1000°C and a stress level of 150 MPa. Fig. 14. (a) Creep strain versus time for Nextel 610/AS tested at 1000°C and four stress levels. In all cases the behavior was essentially linear to failure. (b) Creep strain versus time for Nextel 610/AS tested at 1100°C. For both stress levels the behavior was essentially linear to failure. Table I. Measured Residual Strength of Samples that Reached Run-Out during Fatigue and Creep Rupture Testing Temperature (°C) Test stress level (MPa) Residual strength (MPa)† Modulus (GPa) Fatigue 23 160 228 74 1000 125 214 86 1000 150 225 78 Creep 1000 75 205 76 † All residual strength tests were conducted at room temperature. June 2003 Characterization and High-Temperature Mechanical Behavior of an Oxide/Oxide Composite 987
988 Journal of the American Ceramic Sociery-Zawada et al. Vol 86. No 6 2 Tension Test e 15U Tension Test 101010210310105106107 Creep rupture stress level versus time to failure for Nextel 61O/ AS tested at It000° and 1 100°C L20 Hm Fig. 14(b)is only 25 000 s. When plotted on the same scale 1000C tests, the traces appear to be almost vertical, with MPa test running only "7h. Such short lives at 1100C indicate that this oxide/oxide CMc should be used only 1000.C under static loader Larger creep strain and longer creep life were observed at lower pplied stresses at both temperatures. At steady state, it was assumed that the stress in all phases equilibrated; therefore, the matrix and fiber creep rates should be the same. In most CMCs matrix damage dominates creep rupture lifetimes. At high temper ature, oxidization of the carbon or BN fiber-matrix interphase and the Sic fiber through matrix cracks controls the rupture process and rupture time. 6, 47 Matrix damage already exists in this oxide/ oxide and oxidation is not a concern. The influence of the cracked matrix on creep behavior is not completely understood at this point. Obviously, the fiber axial stresses were far higher than other stress components. Therefore, creep rupture was probably domi- nated by creep rupture of Nextel 610 fibers. The creep rupture of Fig. 16.(a) High-resolution SEM micrograph of a Nextel 610 fiber failed ndividual Nextel 610 fibers was believed to be due to boundar creep rupture. Fracture surface of the fiber is characterized by avity coalescence. but this was often after strains as high as 30% intergranular crack growth. (b) High-resolution SEM micrograph of a accumulated. Postfailure analysis indicated that intergranular ailed creep rupture. Fracture surface clearly shows that matrix remains bonded to the fibers failure was the predominant characteristic failure mode in the ruptured specimens(Fig. 16(a)). The mechanism controlling steady-state creep of Nextel 610 fibers was suggested to be interface-reaction-controlled diffusion creep with fine intergranu- nherently better creep resistance of the N720 fiber. Creep strain lar crack formation. 8-50 The lack of a noticeable tertiary creep versus time traces were similar to the NolO/AS traces recorded in region suggested that there was a spontaneous linkage of creep- this investigation nucleated cracks that occurred abruptly just before failure. Cracks continued to nucleate throughout the creep process until a critica crack density was reached, which caused spontaneous linkage and IV. Discussion failure The 75 MPa test at 1000%C reached run-out and was tested at There are distinct behaviors of porous-matrix oxide room temperature for residual strength. Even though the specimen oxide CMCs that are different from traditional Cmcs with a d experienced almost 1% creep strain, there was no decrease in fiber-matrix interphase. In general, Sic fibers are stronger than oxide fibers. Oxide/oxide CMCs generally exhibit strengths that tensile strength. It was interesting to observe so much creep strain and yet measure no decrease in strength. This suggested that are in the range of 200-250 MPa, whereas Nicalon-containing strength-governing flaws did not enlarge during the initial or CMCs typically range from 200 to 350 MPa for 8HSW cross-ply composites. 8 Even with linear stress versus strain behavior intermediate stages of creep. Observations of the fracture surfaces revealed that most of the exposed fibers on the fracture surtace had relatively notch insensitive during fast fracture.20,22,51,52The natrix remaining on the fibers(Fig. 16(b), which indicated that room-temperature fatigue performance is very similar to many other CMCs, with run-out a high percentage of the average tensile prove the creep resistance of N610/AS strength. However. the 1000%C fatigue limit level of 155 MPa is There is little creep data in the literature for 8HSW oxide/oxide inimical than that measured for most other cmcs composites containing Nextel 610 fibers(N610) with similar fiber volume fractions. A horizontal o/N plot suggests tests were conducted on a SHSW N720/AS al essive 1100C. This composite demonstrated a run-ou SiC-fiber CMCs perform as well in fatigue at 1000.C once the stress level of 150 MPa. Such results clearly documented the matrix is cracked. This is a very important observation, because
Fig. 14(b) is only 25 000 s. When plotted on the same scale as the 1000°C tests, the traces appear to be almost vertical, with the 50 MPa test running only 7 h. Such short lives at 1100°C clearly indicate that this oxide/oxide CMC should be used only below 1000°C under static loading. Larger creep strain and longer creep life were observed at lower applied stresses at both temperatures. At steady state, it was assumed that the stress in all phases equilibrated; therefore, the matrix and fiber creep rates should be the same. In most CMCs, matrix damage dominates creep rupture lifetimes. At high temperature, oxidization of the carbon or BN fiber–matrix interphase and the SiC fiber through matrix cracks controls the rupture process and rupture time.46,47 Matrix damage already exists in this oxide/ oxide, and oxidation is not a concern. The influence of the cracked matrix on creep behavior is not completely understood at this point. Obviously, the fiber axial stresses were far higher than other stress components. Therefore, creep rupture was probably dominated by creep rupture of Nextel 610 fibers. The creep rupture of individual Nextel 610 fibers was believed to be due to boundary cavity coalescence, but this was often after strains as high as 30% accumulated.48 Postfailure analysis indicated that intergranular failure was the predominant characteristic failure mode in the ruptured specimens (Fig. 16(a)). The mechanism controlling steady-state creep of Nextel 610 fibers was suggested to be interface-reaction-controlled diffusion creep with fine intergranular crack formation.48–50 The lack of a noticeable tertiary creep region suggested that there was a spontaneous linkage of creepnucleated cracks that occurred abruptly just before failure. Cracks continued to nucleate throughout the creep process until a critical crack density was reached, which caused spontaneous linkage and failure. The 75 MPa test at 1000°C reached run-out and was tested at room temperature for residual strength. Even though the specimen had experienced almost 1% creep strain, there was no decrease in tensile strength. It was interesting to observe so much creep strain and yet measure no decrease in strength. This suggested that strength-governing flaws did not enlarge during the initial or intermediate stages of creep. Observations of the fracture surfaces revealed that most of the exposed fibers on the fracture surface had matrix remaining on the fibers (Fig. 16(b)), which indicated that load would have been shed to the matrix. A denser matrix would improve the creep resistance of N610/AS. There is little creep data in the literature for 8HSW oxide/oxide composites containing Nextel 610 fibers (N610). Creep rupture tests were conducted on a 8HSW N720/AS composite at 1100°C.51 This composite demonstrated a run-out of 100 h at a stress level of 150 MPa. Such results clearly documented the inherently better creep resistance of the N720 fiber. Creep strain versus time traces were similar to the N610/AS traces recorded in this investigation. IV. Discussion There are several distinct behaviors of porous-matrix oxide/ oxide CMCs that are different from traditional CMCs with a fiber–matrix interphase. In general, SiC fibers are stronger than oxide fibers. Oxide/oxide CMCs generally exhibit strengths that are in the range of 200–250 MPa, whereas Nicalon-containing CMCs typically range from 200 to 350 MPa for 8HSW cross-ply composites.18 Even with linear stress versus strain behavior, oxide/oxide CMCs have shown good fracture toughness and are relatively notch insensitive during fast fracture.20,22,51,52 The room-temperature fatigue performance is very similar to many other CMCs, with run-out a high percentage of the average tensile strength. However, the 1000°C fatigue limit level of 155 MPa is significantly higher than that measured for most other CMCs18 with similar fiber volume fractions. A horizontal /N plot suggests that there is no progressive fatigue damage developing. No SiC-fiber CMCs perform as well in fatigue at 1000°C once the matrix is cracked. This is a very important observation, because Fig. 15. Creep rupture stress level versus time to failure for Nextel 610/AS tested at 1000° and 1100°C. Fig. 16. (a) High-resolution SEM micrograph of a Nextel 610 fiber failed under creep rupture. Fracture surface of the fiber is characterized by intergranular crack growth. (b) High-resolution SEM micrograph of a failed creep rupture. Fracture surface clearly shows that matrix remains bonded to the fibers. 988 Journal of the American Ceramic Society—Zawada et al. Vol. 86, No. 6
June 2003 Characteriation and High-Temperature Mechanical Behavior of an Oxide/Oxide Composite 98 the oxide/oxide CMc is extensively micocracked during process- as a NO radical, may be the cause of strength decrease. However, ing Under service conditions, it is anticipated that cracks eventu- much work remains to determine the reactive species and what ally form in the matrix. In fact, several CMCs have shown concentrations result in loss of fiber strength evidence of matrix cracking at stress levels well below the Degradation of fast-fracture performance because of thermal proportional limit, Significant life must remain in the component exposure is another concern. Although the composite tensile once matrix cracking occurs. Therefore, high tensile strength is not strength decreases only -15% between room temperature and always the most important property. For longer-life applications, 982C, the change in fractography with temperature suggests the CMC has to demonstrate high-temperature durability and damage mechanisms change at higher temperatures. Other studies significant retained strength. N610/AS exhibits excellent retained suggest more drastic decreases in strength at temperature if the strength after reaching run-out, and this behavior makes lifetime composites are notched o One source of change may be from design easier compared with Sic-fiber-containing CMCs degradation of fast-fracture strength of the nextel 610 fiber at high The creep behavior of the oxide/oxide CMC studied is strongly temperatures. High-temperature studies by the manufacturerand governed by the Nextel 610 fiber. Amorphous SiO2 in the matrix others suggest up to a 25% decrease in tensile strength at limits contribution of the matrix to creep resistance. Creep perfor- temperatures as low as 900oC. The change in residual stress state mance in the N610/AS composites can be improved by removing may also impact mechanical properties at high temperature the sio, in the matrix. Oxide fibers that contain mullite. such as Another source of change may be coarsening of porosity in th Nextel 720, exhibit improved creep resistance but lower strength. matrix. However, it is not clear how this changes tensile strengt A combination plot of stress versus time to failure for creep and Composites with Al,O/mullite matrices with much coarser poros- fatigue(Fig. 17)clearly shows shorter life with static loading and ity than N610/AS have similar properties to N610/AS, but the longer life with fatigue. Most Nicalon-fiber-containing CMCs retain better properties after thermal exposure. 63 Another possible show the opposite behavior. 4 Strain accumulation during fatigue high-temperature effect is related to SiOz viscosity. SiO, with is much lower for N610/AS. Cyclic loading does not appear to 0.04%0 12% hydroxyl content has a viscosity of -l X 10-Pa produce measurable damage to the fibers, because the retained at 1100 C37 The 0.05 mm/s displacement rate over a 1.5 cm hot tensile strength is as high or higher than the as-received material zone used in the tension tests causes a stress of -3 GPa in SiO, The residual strength data(Table D) clearly indicate that N61O/As ligaments. This stress may not be sufficient to fracture fine SiO2 is very dependent on the rate at which damage occurs. If strain can ligaments over a short gauge length. Instead they may creep by be accommodated slowly, then there is little flaw growth in the viscous flow under the test conditions used and change the damage nd retained st mechanisms from distributed microcracking ahead of crack tips to The low values of interlaminar strength impact component a mechanism involving viscous flow of SiO, at high temperature design. This property probably cannot be improved without significantly increasing matrix density or using three-dimensional ber architecture The extensive porosity of the matrix raises concems about wear and permeability Staehler and Zawandashowed that a N720/AS le tensile strength and fatigue performance at room CMC tested in an F110 nozzle as a divergent flap experienced emperatures was found in an N610/AS CMC that significant wear on the surface. The wear was a result of chatter weak, porous matrix for a weak fiber-matrix inter impact combined with localized point loading. For locations where ile and compressive strength are moderate at roor the contact was over a wider area, there was no wear. Extensive temperature compared with other CMCs, but, unlike most other orosity also means the CMC is not hermetically sealed, and the CMCs, fatigue performance does not change significantly with interconnecting porosity exposes the fibers to the environment temperature up to 1000oC. The Nextel 610 fiber results in low Zawada has shown that phosphoric acid-containing compounds creep resistance and limits the use time above 1000C. The can produce reactive species during heating to a 1000C that low-modulus, porous, and precracked aluminosilicate matrix ha enetrate the entire composite and decrease strength by >60% in low in-plane and interlaminar strength. The mechanisms by which a mater of minutes. Such rapid loss in strength is also observed mechanical properties degrade above 1000C are problematic during coating of Nextel 610 and Nextel 720 fibers. 57-59 It is Fiber and matrix may be involved uggested that hydrogen ions or nitrogen-containing species, such References IR J Kerans, R S. Hay, N. J. Pagano, and T. A Parthasarathy, "The Role T=1000°c Fiber-Matrix Interface in Ceramic Composites, Amm. Ceram. Soc. Bull., 68[2] Fatigue: R= 0.05, f=1 Hz 429-42(1989) A. G. Evans and F. W. Zok, "Review: The Physics and Mechanics of Fibre- ST. Mah, M. G. Mendiratta, A. P R. Ruh, and K. S. Mazdiyasni Reinforced Glass-Ceramic-Matrix JAm. Cerami.Soc,68[9c-248-C-251(1985) A G. Evans. D. B Ma erization of Glass and Glass-Ceramic matrix/ Nicalon SiC Fiber Composites, "Mater. Sci. Res, 20, 546-60(1986) K. Prewo and J. J. Brennan, "High-Strength Silicon Carbide-Fiber-Reinforced 50 口 Creep RL Glass-Ceramic-Matrix Composites, pp. 387-99 in Ceramic Microstructure tole of Interfaces. Edited by J. A. Pask and A G. Evans. Plenum, New York, 1986 R. F. Cooper and K. Chyung, "Structure and Chemistry of Fiber-Matrix Interfaces 109101102103104105106107 copy Study, "J. Mater. Sci, 22, 3148-60(1987). Interactions, "Ceram. Eng. Sci. Proc., 8[7-8]634-43(1987). 1°R. Naslain,O. Dune, ely, C.R. Brosse, J. P. Rocher, and J. mechanical behavior at 1000%C. Test data are shown for tensile Interphase in Ceramic-Matrix Composites, " J. A Ceram. and creep rupture. For this material, the shortest lives are for Soc,740]2482-88(1991) Nutt, and J. J. Brennan, "Interfacial Microstructure and load, implying this material is substantially more sensitive to cree Chemistry of SiC/BN Dual-Coated Nicalon-Fiber-Reinforced Glass-Ceramic-Matrix Composites, "J. Am. Ceram Soc., 77[5] 1329-39(1994)
the oxide/oxide CMC is extensively micocracked during processing. Under service conditions, it is anticipated that cracks eventually form in the matrix. In fact, several CMCs have shown evidence of matrix cracking at stress levels well below the proportional limit.53 Significant life must remain in the component once matrix cracking occurs. Therefore, high tensile strength is not always the most important property. For longer-life applications, the CMC has to demonstrate high-temperature durability and significant retained strength. N610/AS exhibits excellent retained strength after reaching run-out, and this behavior makes lifetime design easier compared with SiC-fiber-containing CMCs. The creep behavior of the oxide/oxide CMC studied is strongly governed by the Nextel 610 fiber. Amorphous SiO2 in the matrix limits contribution of the matrix to creep resistance. Creep performance in the N610/AS composites can be improved by removing the SiO2 in the matrix. Oxide fibers that contain mullite, such as Nextel 720, exhibit improved creep resistance but lower strength. A combination plot of stress versus time to failure for creep and fatigue (Fig. 17) clearly shows shorter life with static loading and longer life with fatigue. Most Nicalon-fiber-containing CMCs show the opposite behavior.54 Strain accumulation during fatigue is much lower for N610/AS. Cyclic loading does not appear to produce measurable damage to the fibers, because the retained tensile strength is as high or higher than the as-received material. The residual strength data (Table I) clearly indicate that N610/AS is very dependent on the rate at which damage occurs. If strain can be accommodated slowly, then there is little flaw growth in the fibers, and retained strength is high. The low values of interlaminar strength impact component design. This property probably cannot be improved without significantly increasing matrix density or using three-dimensional fiber architecture. The extensive porosity of the matrix raises concerns about wear and permeability. Staehler and Zawanda55 showed that a N720/AS CMC tested in an F110 nozzle as a divergent flap experienced significant wear on the surface. The wear was a result of chatter impact combined with localized point loading. For locations where the contact was over a wider area, there was no wear. Extensive porosity also means the CMC is not hermetically sealed, and the interconnecting porosity exposes the fibers to the environment. Zawada56 has shown that phosphoric acid-containing compounds can produce reactive species during heating to a 1000°C that penetrate the entire composite and decrease strength by 60% in a mater of minutes. Such rapid loss in strength is also observed during coating of Nextel 610 and Nextel 720 fibers.57–59 It is suggested that hydrogen ions or nitrogen-containing species, such as a NO radical, may be the cause of strength decrease. However, much work remains to determine the reactive species and what concentrations result in loss of fiber strength. Degradation of fast-fracture performance because of thermal exposure is another concern. Although the composite tensile strength decreases only 15% between room temperature and 982°C, the change in fractography with temperature suggests damage mechanisms change at higher temperatures. Other studies suggest more drastic decreases in strength at temperature if the composites are notched.60 One source of change may be from degradation of fast-fracture strength of the Nextel 610 fiber at high temperatures. High-temperature studies by the manufacturer61 and others suggest up to a 25% decrease in tensile strength at temperatures as low as 900°C.62 The change in residual stress state may also impact mechanical properties at high temperature. Another source of change may be coarsening of porosity in the matrix. However, it is not clear how this changes tensile strength. Composites with Al2O3/mullite matrices with much coarser porosity than N610/AS have similar properties to N610/AS, but they retain better properties after thermal exposure.63 Another possible high-temperature effect is related to SiO2 viscosity. SiO2 with 0.04%–0.12% hydroxyl content has a viscosity of 1 1012 Pas at 1100°C.37 The 0.05 mm/s displacement rate over a 1.5 cm hot zone used in the tension tests causes a stress of 3 GPa in SiO2 ligaments. This stress may not be sufficient to fracture fine SiO2 ligaments over a short gauge length. Instead they may creep by viscous flow under the test conditions used and change the damage mechanisms from distributed microcracking ahead of crack tips to a mechanism involving viscous flow of SiO2 at high temperature. V. Conclusions Reasonable tensile strength and fatigue performance at room and high temperatures was found in an N610/AS CMC that substitutes a weak, porous matrix for a weak fiber–matrix interface. Tensile and compressive strength are moderate at room temperature compared with other CMCs, but, unlike most other CMCs, fatigue performance does not change significantly with temperature up to 1000°C. The Nextel 610 fiber results in low creep resistance and limits the use time above 1000°C. The low-modulus, porous, and precracked aluminosilicate matrix has low in-plane and interlaminar strength. The mechanisms by which mechanical properties degrade above 1000°C are problematic. Fiber and matrix may be involved. References 1 R. J. Kerans, R. S. Hay, N. J. Pagano, and T. A. Parthasarathy, “The Role of the Fiber–Matrix Interface in Ceramic Composites,” Am. Ceram. Soc. Bull., 68 [2] 429–42 (1989). 2 A. G. Evans and F. W. Zok, “Review: The Physics and Mechanics of FibreReinforced Brittle-Matrix Composites,” J. Mater. Sci., 29, 3857–96 (1994). 3 T. Mah, M. G. Mendiratta, A. P. Katz, R. Ruh, and K. S. Mazdiyasni, “High-Temperature Mechanical Behavior of Fiber-Reinforced Glass-Ceramic-Matrix Composites,” J. Am. Ceram. Soc., 68 [9] C-248–C-251 (1985). 4 H. C. Cao, E. Bischoff, O. Sbaizero, M. Ruhle, A. G. Evans, D. B. Marshall, and J. J. Brennan, “Effect of Interfaces on the Properties of Fiber-Reinforced Ceramics,” J. Am. Ceram. Soc., 73 [6] 1691–99 (1990). 5 J. J. Brennan, “Interfacial Characterization of Glass and Glass-Ceramic Matrix/ Nicalon SiC Fiber Composites,” Mater. Sci. Res., 20, 546–60 (1986). 6 K. Prewo and J. J. Brennan, “High-Strength Silicon Carbide-Fiber-Reinforced Glass-Matrix Composites,” J. Mater. Sci., 15 [2] 463–68 (1980). 7 J. J. Brennan, “Interfacial Chemistry and Bonding in Fiber-Reinforced Glass and Glass-Ceramic-Matrix Composites”; pp. 387–99 in Ceramic Microstructures ’86: Role of Interfaces. Edited by J. A. Pask and A. G. Evans. Plenum, New York, 1986. 8 R. F. Cooper and K. Chyung, “Structure and Chemistry of Fiber–Matrix Interfaces in Silicon Carbide-Fibre-Reinforced Glass-Ceramic Composites: An Electron Microscopy Study,” J. Mater. Sci., 22, 3148–60 (1987). 9 R. N. Singh and M. K. Brun, “Effect of Boron Nitride Coating on Fiber–Matrix Interactions,” Ceram. Eng. Sci. Proc., 8 [7–8] 634–43 (1987). 10R. Naslain, O. Dugne, A. Guette, J. Sevely, C. R. Brosse, J. P. Rocher, and J. Cotteret, “Boron Nitride Interphase in Ceramic-Matrix Composites,” J. Am. Ceram. Soc., 74 [10] 2482–88 (1991). 11E. Y. Sun, S. R. Nutt, and J. J. Brennan, “Interfacial Microstructure and Chemistry of SiC/BN Dual-Coated Nicalon-Fiber-Reinforced Glass-Ceramic-Matrix Composites,” J. Am. Ceram. Soc., 77 [5] 1329–39 (1994). Fig. 17. Combination plot for 8HSW Nextel 610/AS documenting the mechanical behavior at 1000°C. Test data are shown for tensile, fatigue, and creep rupture. For this material, the shortest lives are for sustained load, implying this material is substantially more sensitive to creep than to fatigue. June 2003 Characterization and High-Temperature Mechanical Behavior of an Oxide/Oxide Composite 989
990 Journal of the American Ceramic Sociery-Zawada et al. Vol 86. No 6 12T. Matsuda, "Stability to Moisture for Chemically Vapour Deposited Boron 40A. S. Fareed and G. H. Schiroky, "Microstructure and Properties of Nextel M610 Fiber Reinforced Ceramic and Metal Matrix Composites, Ceram. Eng. Sci Proc., A G Hegedus,. ramie Bodies of controlled Porosity and Process for Making 41344-52(1994) J. A. Heathcote, X.-Y. Gong, J. Y. Yang, U. Ramamurty, and F. w. Zo arrison,"Fiber-Reinforced Ceramic cal Properties of an All-Oxide Ceramic Composite, J. Am. Ceram. Composite Member, U.S. Pat. No. 5488017, Jan 30, 1996. Soc,820]2721-30(1999 A. Szweda, M. L. Millard, and M. G. Harrison, ""Fiber-Reinforced Ceramic Matrix Composite Member and Method for Making, U.S. Pat. No 5601 674, Feb. Composites, " Thermal and Mechanical Test Methods and Behavior of Continuous J. Dunvak, D. R. Chang, and M. L. Millard. "Thermal Aging ef opposites, ASTM, STP 1309, 31-48(1997) Oxide/Oxide Ceramic-Matrix Composites", pp. 675-90 in /7th Conference Composites, "Acta Metall. Mater, 43 [3]859-75(1995). 3235, Part 2. Edited by J. D. Buckley. NAS J. w. Holmes and C Cho,"Experimental Observations of Frictional Heating in a Edited by w.S. Coblenz. ARPA Ceramic Technology Insertion Program(DARPA) osite, M.S. Thesis, Air Force Institute of Technology, Annapolis, MD, 1994 L. P. Zawada and S. S. Lee. "Evaluation of the Fatigue Performance of Five J. w. Holmes and x. Wu, "Elevated space Applications", pp. 1669-74 in Sixth International Fatigue Fiber-Reinforced Ceramics", pp. 193-260 in Elerated Temperature Creep Behavior Congress, VoL. Ill. Edited by G. Lutjering and H. Nowack. Elsevier Science cs. Edited by S. V. Nair and K. Jakus NY,1996 IT. J. Lu, "Crack Branching in All-Oxide Ceramic Composites, "J.Am.Ceram 4D. W. Meyer, R F. Cooper, and M E Plesha, "High-Temperature Creep and the w.-C. Tu, F. F, Lange, and A. G. Evans, "Concept for a Damage-Tolerant Ceramic Composite with"Strong'Interfaces, " J. Am. Ceram Soc., 79[2]417-24 4D. J. Pysher and R. E. Tressler, "Creep Rupture Studies of Two Alumina-Based 2D. M. Carper and M. L. Millard, Oxide-Oxide CMCs for Engine Exhaust 3213(92 Systems, HAVE FORM, U.S. Air Force, Fort Walton Beach, FL, 19 stalline Oxide Fibers, Ceram. Eng Sai. Proc., 13[7-8] 218-26 -F. W. Zok and C.G. Ley ies of Porous- Matrix Ceramic Composites,Ad. Eng. Mater, 3[1-2] 15-23(2001) SOD. M. Wilson, D. C Lunenburg, and S. L. Lieder, "High-Temperature Properties M. K. Cinibulk, J.R. Welch, and R. S. Hay, ""Preparation of Thin Sections of of Nextel 610 and Alumina-Based Nanocomposite Fibers,Ceram. En g. Sci. Proc. Coated Fibers for Character by Transmission Electron Microscopy, J.Am SR. John, D J. Buchanan, and L. P Zawada, "Creep Deformation and Rupture 2G. A. Hartman and N. E. Ashbaugh,A Mechanics Test Automation Behavior of a Notched Oxide/Oxide Nextel 720/AS Composite, Ceram. Eng Sc n Basic Research Lab Pmoc,2113567-73(2000) igue Test Methodology and Matrix Composites at ASTA. STP 1157. 52-68 (199 and Elevated Temperature Oxide/Oxide Ceramic-Matrix Composite, Mechanical, Thermal, and Ermvironmente Testing and Performance of Ceramic Composites and Components, ASTM, STP zR. K. Bordia and A Jagota."Crack Growth and Damage in Constrained Sintering Films,J Am Ceram. Soc, 76[10]2475-85(1993). J Z Gyekenyesi, and R. T. Bhatt, ""Damage Accumulation 2C. P. Ostertag, P. G. Charalambrides, and A. G. Evans, "Observations and 2-D Woven SiC/SiC Ceramic-Matrix Composites, Mechanical, Thermal, an Analysis of Sintering Damage, Acta Metall, 37 [71 2077-84(1989) ASTM STE1392,306-19(2000 Microstructural Development in an Al,O, Matrix Containing a Large Volume Fraction of ZrO2 Inclusions, "J.Anm Ceram. Soc., 75 3]519-24(19 s4S. S. Lee, L. P. Zawada, J. M. Stachler, and C. A Folsom, "Mechanical Behavior 2D. C. C. Lam and FF. Lange.""Microstructual Observations on Constrained and High-Temperature Performance of a Woven Nicalon/SiNC Ceramic-Matrix JAm. Ceram.Soc,811797-11(1998 sSJ. M. Staehler and L. P. Zawada, "Performance of Four Ceramic-Matrix sG. W. Scherer, "Sintering with Rigid Inclusions, "J. Am. Ceram Soc., 70 [10] Composite Divergent Flap Inserts Following Ground Testing on an F110 Turbofa Engine,JAm Ceram. Soc., 83[7 1727-38(2000). J. A. Pask and A P. Tomsia, "Formation of Mullite from Sol-Gel Mixtures and ile behavior of Several Kaolinite,".A. Ceram Soc., 74 [10]2367-73(1991) Oxide/Oxide Composites, Ceram. Eng. Sci Proc., 19[ 3]327-39( 32D. X. Li and W. J. Thomson, on from nonstoichiometric E. Boakye, R S Hay, and M. D. Petry, "Continuous Coating of Oxide Fiber Diphasic Precursors,J Am Ceram Soc., 74 (10 2382-87(1991). Tows Using Liquid Precursors: Monazite Coatings on Nextel 720, "J. Am. Ceram. G. w. Scherer, ""Viscous Sintering of a Bimodal Pore-Size Distribution,"J.Am. Soc,821912321-31(999 Ceram.Soc,67[l709-15(1984 SsR. S. Hay, E. E. Boakye, and M. D. Petry, "Effect of Coating Inclusions,"JAm. Ceram. Soc., 71 Togi with a Pore-Size Distribution and Rigid m.Soc,20,589- w. Scherer, "Coarsening in a Viscous Matrix, "J.A. Ceram Soc., 81 [11 49-54(1998) 2793801(2001) SN. M. Parikh, "Effect of Atmosphere on Surface Tension of Glass,"J.Am bov. A. Kramb, R John, and L. P. Zawada, "Notched Fracture Behavior Oxide/Oxide Ceramic-Matrix Composite, J. Am. Ceram. Soc., 82 [11] 3087-96 6D. M. Wilson and L. R Visser. "High-Performance Oxide Fibers for Metal and M. N, Rahaman,"Theory of Solid-State and Viscous Sintering": Ch. 8, pp Ceramic Composites, Composites, Part 4, 32, 1143-53(200 374-444 in Ceramic Processing and Sintering, Ist ed Marcel-Dekker, New York, 6 Z R. Xu, K.K. Chawla, and X Li, "Effect of High-Temperature Exposure on the Tensile Strength of Alumina Fiber Nextel 610, Mater. Sci. Eng. A, 171, 249-56 J. C. Aveston, C. G. Cooper, and A. Kelly, "Single and Multiple Fracture"; pp 15-26 in Properties of Fiber Composites, Proeedings of National Physics Laboratory G. Levi, J. Y. Yang, B.J. Dalgleish, F. w. Zok, and A G. Evans, "Processing Conference, Edited by J. C. Aveston, C. G. Cooper, and A. Kelly. IPC Science and and Performance of an All-Oxide Ceramic Composite, "J.Amm. Ceram. Soc., &1 [8 Technology Press, Guilford, U. K, 197 2077-86(1998
12T. Matsuda, “Stability to Moisture for Chemically Vapour Deposited Boron Nitride,” J. Mater. Sci., 24, 2353–58 (1989). 13A. G. Hegedus, “Ceramic Bodies of Controlled Porosity and Process for Making Same,” U.S. Pat. No. 5 017 522, May 21, 1991. 14A. Szweda, M. L. Millard, and M. G. Harrison, “Fiber-Reinforced Ceramic Composite Member,” U.S. Pat. No. 5 488 017, Jan. 30, 1996. 15A. Szweda, M. L. Millard, and M. G. Harrison, “Fiber-Reinforced CeramicMatrix Composite Member and Method for Making,” U.S. Pat. No. 5 601 674, Feb. 11, 1997. 16T. J. Dunyak, D. R. Chang, and M. L. Millard, “Thermal Aging Effects in Oxide/Oxide Ceramic-Matrix Composites”; pp. 675–90 in 17th Conference on Metal Matrix, Carbon, and Ceramic Matrix Composites, NASA Conference Publication 3235, Part 2. Edited by J. D. Buckley. NASA, Washington, DC, 1993. 17L. P. Zawada and S. S. Lee, “Mechanical Behavior of CMCs for Flaps and Seals”; pp. 267–322 in Defense Advanced Projects Agency Conference Publication. Edited by W. S. Coblenz. ARPA Ceramic Technology Insertion Program (DARPA), Annapolis, MD, 1994. 18L. P. Zawada and S. S. Lee, “Evaluation of the Fatigue Performance of Five CMCs for Aerospace Applications”; pp. 1669–74 in Sixth International Fatigue Congress, Vol. III. Edited by G. Lutjering and H. Nowack. Elsevier Science, Tarrytown, NY, 1996. 19T. J. Lu, “Crack Branching in All-Oxide Ceramic Composites,” J. Am. Ceram. Soc., 79 [1] 266–74 (1996). 20W.-C. Tu, F. F. Lange, and A. G. Evans, “Concept for a Damage-Tolerant Ceramic Composite with ‘Strong’ Interfaces,” J. Am. Ceram. Soc., 79 [2] 417–24 (1996). 21D. M. Carper and M. L. Millard, “Oxide–Oxide CMCs for Engine Exhaust Systems,” HAVE FORM, U.S. Air Force, Fort Walton Beach, FL, 1997. 22F. W. Zok and C. G. Levi, “Mechanical Properties of Porous-Matrix Ceramic Composites,” Adv. Eng. Mater., 3 [1–2] 15–23 (2001). 23M. K. Cinibulk, J. R. Welch, and R. S. Hay, “Preparation of Thin Sections of Coated Fibers for Characterization by Transmission Electron Microscopy,” J. Am. Ceram. Soc., 79 [9] 2481–84 (1996). 24G. A. Hartman and N. E. Ashbaugh, “A Fracture Mechanics Test Automation System for a Basic Research Laboratory,” ASTM, STP 1092, 95–110 (1990). 25L. M. Butkus, L. P. Zawada, and G. A. Hartman, “Fatigue Test Methodology and Results for Ceramic-Matrix Composites at Room and Elevated Temperatures,” ASTM, STP 1157, 52–68 (1990). 26R. K. Bordia and A. Jagota, “Crack Growth and Damage in Constrained Sintering Films,” J. Am. Ceram. Soc., 76 [10] 2475–85 (1993). 27C. P. Ostertag, P. G. Charalambrides, and A. G. Evans, “Observations and Analysis of Sintering Damage,” Acta Metall., 37 [7] 2077–84 (1989). 28O. Sudre and F. F. Lange, “The Effect of Inclusions on Densification: I, Microstructural Development in an Al2O3 Matrix Containing a Large Volume Fraction of ZrO2 Inclusions,” J. Am. Ceram. Soc., 75 [3] 519–24 (1992). 29D. C. C. Lam and F. F. Lange, “Microstructual Observations on Constrained Densification of Alumina Powder Containing a Periodic Array of Sapphire Fibers,” J. Am. Ceram. Soc., 77 [7] 1976–78 (1994). 30G. W. Scherer, “Sintering with Rigid Inclusions,” J. Am. Ceram. Soc., 70 [10] 719–25 (1987). 31J. A. Pask and A. P. Tomsia, “Formation of Mullite from Sol–Gel Mixtures and Kaolinite,” J. Am. Ceram. Soc., 74 [10] 2367–73 (1991). 32D. X. Li and W. J. Thomson, “Mullite Formation from Nonstoichiometric Diphasic Precursors,” J. Am. Ceram. Soc., 74 [10] 2382–87 (1991). 33G. W. Scherer, “Viscous Sintering of a Bimodal Pore-Size Distribution,” J. Am. Ceram. Soc., 67 [11] 709–15 (1984). 34G. W. Scherer, “Viscous Sintering with a Pore-Size Distribution and Rigid Inclusions,” J. Am. Ceram. Soc., 71 [10] C-447–C-448 (1988). 35G. W. Scherer, “Coarsening in a Viscous Matrix,” J. Am. Ceram. Soc., 81 [1] 49–54 (1998). 36N. M. Parikh, “Effect of Atmosphere on Surface Tension of Glass,” J. Am. Ceram. Soc., 41 [1] 18–22 (1958). 37G. Hetherington, K. H. Jack, and J. C. Kennedy, “Viscosity of Vitreous Silica,” Phys. Chem. Glasses, 5 [5] 130–36 (1964). 38M. N. Rahaman, “Theory of Solid-State and Viscous Sintering”; Ch. 8, pp. 374–444 in Ceramic Processing and Sintering, 1st ed. Marcel-Dekker, New York, 1995. 39J. C. Aveston, C. G. Cooper, and A. Kelly, “Single and Multiple Fracture”; pp. 15–26 in Properties of Fiber Composites, Proeedings of National Physics Laboratory Conference. Edited by J. C. Aveston, C. G. Cooper, and A. Kelly. IPC Science and Technology Press, Guilford, U.K., 1971. 40A. S. Fareed and G. H. Schiroky, “Microstructure and Properties of NextelTM 610 Fiber Reinforced Ceramic and Metal Matrix Composites,” Ceram. Eng. Sci. Proc., 15 [4] 344–52 (1994). 41J. A. Heathcote, X.-Y. Gong, J. Y. Yang, U. Ramamurty, and F. W. Zok, “In-Plane Mechanical Properties of an All-Oxide Ceramic Composite,” J. Am. Ceram. Soc., 82 [10] 2721–30 (1999). 42E. Lara-Curzio and M. K. Ferber, “Shear Strength of Continuous-Fiber Ceramic Composites,” Thermal and Mechanical Test Methods and Behavior of Continuous Fiber Ceramic Composites, ASTM, STP 1309, 31–48 (1997). 43A. G. Evans, F. W. Zok, and R. M. McMeeking, “Fatigue of Ceramic-Matrix Composites,” Acta Metall. Mater., 43 [3] 859–75 (1995). 44J. W. Holmes and C. Cho, “Experimental Observations of Frictional Heating in a Fiber-Reinforced Ceramic,” J. Am. Ceram. Soc., 75 [4] 928–38 (1992). 45S. G. Steel, “Monotonic and Fatigue Loading Behavior of an Oxide/Oxide Ceramic-Matrix Composite”; M.S. Thesis. Air Force Institute of Technology, Wright-Patterson Air Force Base, OH, 2000. 46J. W. Holmes and X. Wu, “Elevated Temperature Creep Behavior of ContinuousFiber-Reinforced Ceramics”; pp. 193–260 in Elevated Temperature Creep Behavior of Continuous Fiber-Reinforced Ceramics. Edited by S. V. Nair and K. Jakus. Butterworth-Hienneman, Kent, U.K., 1995. 47D. W. Meyer, R. F. Cooper, and M. E. Plesha, “High-Temperature Creep and the Interfacial Mechanical Response of a Ceramic-Matrix Composite,” Acta Metall. Mater., 41 [11] 3157–70 (1993). 48D. J. Pysher and R. E. Tressler, “Creep Rupture Studies of Two Alumina-Based Ceramic Fibers,” J. Mater. Sci., 27, 423–28 (1992). 49D. J. Pysher and R. E. Tressler, “Tensile Creep Rupture Behavior of AluminaBased Polycrystalline Oxide Fibers,” Ceram. Eng. Sci. Proc., 13 [7–8] 218–26 (1992). 50D. M. Wilson, D. C. Lunenburg, and S. L. Lieder, “High-Temperature Properties of Nextel 610 and Alumina-Based Nanocomposite Fibers,” Ceram. Eng. Sci. Proc., 14 [7–8] 609–21 (1993). 51R. John, D. J. Buchanan, and L. P. Zawada, “Creep Deformation and Rupture Behavior of a Notched Oxide/Oxide Nextel 720/AS Composite,” Ceram. Eng. Sci. Proc., 21 [3] 567–73 (2000). 52R. John, D. J. Buchannan, and L. P. Zawada, “Notch-Sensitivity of a Woven Oxide/Oxide Ceramic-Matrix Composite,” Mechanical, Thermal, and Environmental Testing and Performance of Ceramic Composites and Components, ASTM, STP 1392, 160–71 (2000). 53G. N. Morscher, J. Z. Gyekenyesi, and R. T. Bhatt, “Damage Accumulation in 2-D Woven SiC/SiC Ceramic-Matrix Composites,” Mechanical, Thermal, and Environmental Testing and Performance of Ceramic Composites and Components, ASTM, STP 1392, 306–19 (2000). 54S. S. Lee, L. P. Zawada, J. M. Staehler, and C. A. Folsom, “Mechanical Behavior and High-Temperature Performance of a Woven NicalonTM/SiNC Ceramic-Matrix Composite,” J. Am. Ceram. Soc., 81 [7] 1797–11 (1998). 55J. M. Staehler and L. P. Zawada, “Performance of Four Ceramic-Matrix Composite Divergent Flap Inserts Following Ground Testing on an F110 Turbofan Engine,” J. Am. Ceram. Soc., 83 [7] 1727–38 (2000). 56L. P. Zawada, “Longitudinal and Transthickness Tensile Behavior of Several Oxide/Oxide Composites,” Ceram. Eng. Sci. Proc., 19 [3] 327–39 (1998). 57E. Boakye, R. S. Hay, and M. D. Petry, “Continuous Coating of Oxide Fiber Tows Using Liquid Precursors: Monazite Coatings on Nextel 720,” J. Am. Ceram. Soc., 82 [9] 2321–31 (1999). 58R. S. Hay, E. E. Boakye, and M. D. Petry, “Effect of Coating Deposition Temperature on Monazite Coated Fiber,” J. Eur. Ceram. Soc., 20, 589–97 (2000). 59E. E. Boakye, R. S. Hay, P. Mogilevsky, and L. M. Douglas, “Monazite Coatings on Fibers: II, Coating without Strength Degradation,” J. Am. Ceram. Soc., 84 [12] 2793–801 (2001). 60V. A. Kramb, R. John, and L. P. Zawada, “Notched Fracture Behavior of an Oxide/Oxide Ceramic-Matrix Composite,” J. Am. Ceram. Soc., 82 [11] 3087–96 (1999). 61D. M. Wilson and L. R. Visser, “High-Performance Oxide Fibers for Metal and Ceramic Composites,” Composites, Part A, 32, 1143–53 (2001). 62Z. R. Xu, K. K. Chawla, and X. Li, “Effect of High-Temperature Exposure on the Tensile Strength of Alumina Fiber Nextel 610,” Mater. Sci. Eng. A, 171, 249–56 (1993). 63C. G. Levi, J. Y. Yang, B. J. Dalgleish, F. W. Zok, and A. G. Evans, “Processing and Performance of an All-Oxide Ceramic Composite,” J. Am. Ceram. Soc., 81 [8] 2077–86 (1998). 990 Journal of the American Ceramic Society—Zawada et al. Vol. 86, No. 6