《复合材料 Composites》课程教学资源（学习资料）第五章 陶瓷基复合材料_Mechanical properties of Hi-Nicalon fiber-reinforced celsian

Availableonlineatwww.sciencedirect.com Science Direct E噩≈RS ELSEVIER Journal of the European Ceramic Society 29(2009)525-535 www.elsevier.comlocate/jeurceramsoc Mechanical properties of Hi-Nicalon fiber-reinforced celsian composites after high-temperature exposures in air Narottam p bansal Structures and Materials Division, NASA Glenn Research Center. Cleveland OH 44135 USA Received 21 April 2008: received in revised form 13 June 2008; accepted 19 June 2008 Available online 26 July 2008 Abstract BN/SiC-coated Hi-Nicalon fiber-reinforced celsian matrix composites(CMCs) were annealed for 100 h in air at various temperatures to 1200C, followed by flexural strength measurements at room temperature. Values of yield stress and strain, ultimate strength, and composite modulus remain almost unchanged for samples annealed up to 1100C. a thin porous layer formed on the surface of the 1100 C annealed sample and its density decreased from 3.09 to 2.90 g/em. The specimen annealed at 1200C gained 0.43% weight, was severely deformed, and was covered with a porous layer of thick shiny glaze which could be easily peeled off. Some gas bubbles were also present on the surface. This surface layer consisted of elongated crystals of monoclinic celsian and some amorphous phase(s). The fibers in this surface ply of the CMC had broken into small pieces. The fiber-matrix interface strength was characterized through fiber push-in technique. Values of debond stress, ad, and frictional sliding stress, tf, for the as-fabricated CMC were 0.31+0. 14 GPa and 10.4+3. 1 MPa, respectively. These values compared with 0.53+0.47 GPa and 8.33+ 1.72 MPa for the fibers in the interior of the 1200C annealed sample, indicating hardly any change in fiber-matrix interface strength. The effects of thermal aging on microstructure were investigated using scanning electron microscopy. Only the surface ply of the 1200C annealed specimens had degraded from oxidation whereas the bulk interior part of the CMC was unaffected. A mechanism is proposed explaining the various steps involved during the degradation of the CMC on annealing in air at 1200C Published by elsevier Ltd Keywords: Ceramic composites; Mechanical properties; SiC fibers; Barium aluminosilicate: Fiber-matrix interface 1. Introduction hot sections of turbine engines Results for Nicalon and hi- Nicalon fiber-reinforced celsian matrix composites have been Fiber-reinforced ceramic matrix composites(CMCs)are reported earlier. 6-17 During high-temperature use, CMC com- prospective candidate materials for high gh-temperature struct ponents are prone to degradation in their mechanical properties applications in various industries such as aerospace, power due to oxidation. Tensile, flexural, and shear properties, at generation, energy conservation, nuclear, petrochemical, and temperatures up to 1200C in air, have been reported for cel transportation. A number of ceramic and glass-ceramic com- sian matrix composites reinforced with Nicalon as well as posite systems. are being developed in various research Hi-Nicalon2 14.16 fibers. However, no information is available laboratories. Barium aluminosilicate with monoclinic celsian about the influence of long-term high-temperature exposures on phase is one of the most refractory glass-ceramics. It has a the mechanical properties of these CMCs. The primary objective melting point of >1700C, is phase stable to 1600C, and of this study was to investigate the effects of high-temperature is oxidation resistant. Over the last few years, at NASA Glenn annealing in oxidizing environment on the mechanical proper Research Center, celsian matrix composites'-9reinforced with ties and microstructural stability of Hi-Nicalon fiber-reinforced silicon carbide-based fibers have been investigated for use in celsian matrix composites. The room temperature strength of the composites, after annealing in air at various temperatures from 550 to 1200C, was measured in three-point flexure. The Tel:+12164333855 fiber-matrix interface strength was analyzed using a fiber push E-mail address: narottamP bansal(@ nasa.gov. in technique. 5 0955-2219/S-see front matter. Published by Elsevier Ltd. doi: 10.1016/j-jeurceramsoc200806.023

Available online at www.sciencedirect.com Journal of the European Ceramic Society 29 (2009) 525–535 Mechanical properties of Hi-Nicalon fiber-reinforced celsian composites after high-temperature exposures in air Narottam P. Bansal ∗ Structures and Materials Division, NASA Glenn Research Center, Cleveland, OH 44135, USA Received 21 April 2008; received in revised form 13 June 2008; accepted 19 June 2008 Available online 26 July 2008 Abstract BN/SiC-coated Hi-Nicalon fiber-reinforced celsian matrix composites (CMCs) were annealed for 100 h in air at various temperatures to 1200 ◦C, followed by flexural strength measurements at room temperature. Values of yield stress and strain, ultimate strength, and composite modulus remain almost unchanged for samples annealed up to 1100 ◦C. A thin porous layer formed on the surface of the 1100 ◦C annealed sample and its density decreased from 3.09 to 2.90 g/cm3. The specimen annealed at 1200 ◦C gained 0.43% weight, was severely deformed, and was covered with a porous layer of thick shiny glaze which could be easily peeled off. Some gas bubbles were also present on the surface. This surface layer consisted of elongated crystals of monoclinic celsian and some amorphous phase(s). The fibers in this surface ply of the CMC had broken into small pieces. The fiber–matrix interface strength was characterized through fiber push-in technique. Values of debond stress, σd, and frictional sliding stress, τf, for the as-fabricated CMC were 0.31 ± 0.14 GPa and 10.4 ± 3.1 MPa, respectively. These values compared with 0.53 ± 0.47 GPa and 8.33 ± 1.72 MPa for the fibers in the interior of the 1200 ◦C annealed sample, indicating hardly any change in fiber–matrix interface strength. The effects of thermal aging on microstructure were investigated using scanning electron microscopy. Only the surface ply of the 1200 ◦C annealed specimens had degraded from oxidation whereas the bulk interior part of the CMC was unaffected. A mechanism is proposed explaining the various steps involved during the degradation of the CMC on annealing in air at 1200 ◦C. Published by Elsevier Ltd. Keywords: Ceramic composites; Mechanical properties; SiC fibers; Barium aluminosilicate; Fiber–matrix interface 1. Introduction Fiber-reinforced ceramic matrix composites (CMCs) are prospective candidate materials for high-temperature structural applications in various industries such as aerospace, power generation, energy conservation, nuclear, petrochemical, and transportation. A number of ceramic and glass–ceramic com￾posite systems1,2 are being developed in various research laboratories. Barium aluminosilicate with monoclinic celsian phase is one of the most refractory glass–ceramics. It has a melting point of >1700 ◦C, is phase stable to ∼1600 ◦C, and is oxidation resistant. Over the last few years, at NASA Glenn Research Center, celsian matrix composites3–9 reinforced with silicon carbide-based fibers have been investigated for use in ∗ Tel.: +1 216 433 3855. E-mail address: narottam.p.bansal@nasa.gov. hot sections of turbine engines. Results for Nicalon and Hi￾Nicalon fiber-reinforced celsian matrix composites have been reported earlier.6–17 During high-temperature use, CMC com￾ponents are prone to degradation in their mechanical properties due to oxidation. Tensile, flexural, and shear properties, at temperatures up to 1200 ◦C in air, have been reported for cel￾sian matrix composites reinforced with Nicalon6 as well as Hi-Nicalon12,14,16 fibers. However, no information is available about the influence of long-term high-temperature exposures on the mechanical properties of these CMCs. The primary objective of this study was to investigate the effects of high-temperature annealing in oxidizing environment on the mechanical proper￾ties and microstructural stability of Hi-Nicalon fiber-reinforced celsian matrix composites. The room temperature strength of the composites, after annealing in air at various temperatures from 550 to 1200 ◦C, was measured in three-point flexure. The fiber–matrix interface strength was analyzed using a fiber push￾in technique.15 0955-2219/$ – see front matter. Published by Elsevier Ltd. doi:10.1016/j.jeurceramsoc.2008.06.023

N P Bansal/ Joumal of the European Ceramic Society 29(2009)525-535 barrier to diffusion of boron from bn into the oxide matrix and also prevents diffusion of matrix elements into the fiber. c8Eo 9000000 The matrix of 0.75 Ba0-0.25Sr0-Al2O3-2SiO2(BSAs)com- position was synthesized by a solid-state reaction method as described earlier. 0 The advantage of BSAS over BAS as matrix has been explained earlier. 1.20 Briefly speaking, hexacelsian is the first phase to form in both BAS and SAS systems On heat treatment at -1200C or higher temperatures, transformation of hexacelsian to monoclinic celsian phase is very sluggish in 20 BAS and very rapid in SAs. However, it is known that substi 10 tution of about 25 mol% of Bao with Sro in bAs accelerates the transformation of hexacelsian to the desired monoclinic celsian phase. The experimental setup and the procedure used for fabrica g 1. Scanning Auger microprobe depth profiles of various elements for Hi- Nicalon fibers having a duplex "BN/SiC surface coating deposited by CVD tion of the fiber-reinforced celsian matrix CMC were essentially the same as described earlier. The matrix precursor powder was made into a slurry by dispersing in an organic solvent along 2. Materials and experimental methods with organic additives as binder. surfactant, deflocculant and plasticizer followed by ball milling. Tows of BN/SiC-coated Polymer derived Hi-Nicalon fiber tows(1800 denier, 500 Hi-Nicalon fibers were coated with the matrix precursor by filaments/tow) with low oxygen content from Nippon Carbon passing through the slurry and winding on a rotating drum. Co. were used as the reinforcement. 8, 19 A duplex surface After drying, the prepreg tape was cut to size. Unidirectional layer of boron nitride(Bn) over coated with silicon carbide fiber-reinforced composites were prepared by tape lay-up(12 was applied on the fibers by a commercial vendor using a plies) followed by warm pressing to form a"green"com- continuous chemical vapor deposition(CVD)reactor. The Bn posite. The fugitive organics were slowly burned out of the ating was deposited at 1000C utilizing a proprietary pre- sample in air, followed by hot pressing under vacuum in a cursor and was amorphous to partly turbostratic in nature. a graphite die to yield dense composites. The oxide precursor thin overcoating of SiC was also deposited by CVD onto the was converted into the desired monoclinic celsian phase in situ BN-coated fibers. The SiC layer was crystalline. The nomi- during hot pressing as was confirmed from X-ray diffraction nal coating thicknesses were 0.4 um for BN, and 0.3 um for The hot pressed CMC panel -ll.I cm x 5cm(4.5 in. x 2 in. SiC. The Bn interfacial layer was intended to be a weak, was annealed in argon at 1100C for 2 h and machined into crack deflecting phase, while the SiC overcoat was used as a test bars(50 mm x 0.625 mm x 2.4 mm) for high-temperature ·b 回沙100m m Fig. 2. SEM micrographs at different magnifications showing polished cross-section of a unidirectional Hi-Nicalon/BNSIC/BSAS composite

526 N.P. Bansal / Journal of the European Ceramic Society 29 (2009) 525–535 Fig. 1. Scanning Auger microprobe depth profiles of various elements for Hi￾Nicalon fibers having a duplex “BN/SiC” surface coating deposited by CVD. 2. Materials and experimental methods Polymer derived Hi-Nicalon fiber tows (1800 denier, 500 filaments/tow) with low oxygen content from Nippon Carbon Co. were used as the reinforcement.18,19 A duplex surface layer of boron nitride (BN) over coated with silicon carbide was applied on the fibers by a commercial vendor using a continuous chemical vapor deposition (CVD) reactor. The BN coating was deposited at ∼1000 ◦C utilizing a proprietary pre￾cursor and was amorphous to partly turbostratic in nature. A thin overcoating of SiC was also deposited by CVD onto the BN-coated fibers. The SiC layer was crystalline. The nomi￾nal coating thicknesses were 0.4 m for BN, and 0.3 m for SiC. The BN interfacial layer was intended to be a weak, crack deflecting phase, while the SiC overcoat was used as a barrier to diffusion of boron from BN into the oxide matrix and also prevents diffusion of matrix elements into the fiber. The matrix of 0.75BaO–0.25SrO–Al2O3–2SiO2 (BSAS) com￾position was synthesized by a solid-state reaction method as described earlier.20 The advantage of BSAS over BAS as matrix has been explained earlier.11,20 Briefly speaking, hexacelsian is the first phase to form in both BAS and SAS systems. On heat treatment at ∼1200 ◦C or higher temperatures, transformation of hexacelsian to monoclinic celsian phase is very sluggish in BAS and very rapid in SAS.8 However, it is known that substi￾tution of about 25 mol% of BaO with SrO in BAS accelerates the transformation23 of hexacelsian to the desired monoclinic celsian phase. The experimental setup and the procedure used for fabrica￾tion of the fiber-reinforced celsian matrix CMC were essentially the same as described earlier.10,11 The matrix precursor powder was made into a slurry by dispersing in an organic solvent along with organic additives as binder, surfactant, deflocculant and plasticizer followed by ball milling. Tows of BN/SiC-coated Hi-Nicalon fibers were coated with the matrix precursor by passing through the slurry and winding on a rotating drum. After drying, the prepreg tape was cut to size. Unidirectional fiber-reinforced composites were prepared by tape lay-up (12 plies) followed by warm pressing to form a “green” com￾posite. The fugitive organics were slowly burned out of the sample in air, followed by hot pressing under vacuum in a graphite die to yield dense composites. The oxide precursor was converted into the desired monoclinic celsian phase in situ during hot pressing as was confirmed from X-ray diffraction. The hot pressed CMC panel ∼11.1 cm × 5 cm (4.5 in. × 2 in.) was annealed in argon at 1100 ◦C for 2 h and machined into test bars (∼50 mm × 0.625 mm × 2.4 mm) for high-temperature Fig. 2. SEM micrographs at different magnifications showing polished cross-section of a unidirectional Hi-Nicalon/BN/SiC/BSAS composite.

N P Bansal /Journal of the European Ceramic Sociery 29(2009)525-535 X-ray diffraction(XRD) patterns were recorded at room tem- 2. Hi-Nicalon/BN/SIC/BSAS CMC rature using a step scan procedure(0.02%/20 step, time/step v=043)#1-29-96 0.5 or I s)on a Phillips ADP-3600 automated diffractometer BN/SIC-coated Hi-Nicalon fiber equipped with a crystal monochromator employing Cu Ko radi- measured from dimensions as by the archimedes method Microstructures of the polished cross-sections and fracture surfaces were observed in a JEOL JSM-840A scanning electron microscope. Prior to analysis, a thin layer of carbon was evaporated onto the SEM specimens for electrical conductivity. The elemental compositions of the fibersurface coatings were TGA curves for BSAS monolith, BN/SiC-coated Hi-Nicalon fiber and analyzed with a scanning Auger microprobe(Fisons Instru- alon/BN/SC/BSAS composite recorded at a heating rate of 5.C/min in ments Microlab Model 310-F) The fibers for this analysis were mounted on a stainless steel sample mount by tacking the ends with colloidal graphite. Depth profiling was per- exposures in air and mechanical testing. The volume fraction of formed by sequential ion-beam sputtering and Auger analysis fibers in the composite was found to be -0.32. The ion etching was done with 3 keV argon ions rastered For high-temperature annealing, the CMC bars were rested over an approximately I mm2 area. The etch rate in Ta2O5 on the edges of an alumina boat placed inside a programmable under these conditions was 0.05 nm/s. Auger electron spec- box furnace. The furnace temperature was raised at a heating troscopy(AES)analysis of the coated Hi-Nicalon fibers was rate of 20C/min. CMC bars were annealed at 550, 800, 900, performed using an electron beam current of approximately 1000, 1100, and 1200C for 100 h in stagnant ambient air and 1.5 nA. The beam was rastered over a 2 um x 20 um area of furnace cooled. Dimensions and weight of each test bar were the fiber with the long axis of the area aligned with the long recorded before and after annealing fiber axis. Spectra were acquired in integral mode at beam Mechanical properties were determined from apparent energy of 2 keV and depth profiles were generated by plot tress-strain curves recorded from a three-point flexure test ting elemental peak areas against ion etch time. The atomic speed of 1.27 mm/min(0.05 in /min) and support span(L)of the spectrometer transmsionrr dividing the peak areas by flexure test bars. Stress, o, was calculated from beam theory, sensitivity factors were derived from spectra of ion etched Si, assuming a linear elastic beam, using the equation B, SiC, BN, and T102 standards. The depth scale is from the Ta2O5 calibration and no attempt has been made to adjust (1) for the actual etch rate for each material. Only the fibers with a smooth surface coating. rather than those having thick where b and h are the width and thickness of the test sample and and rough coating morphologies, were used for Auger analy- P is the load. The yield stress, y, was taken from the onset fsis deviation from linearity in the stress-strain curve. Elastic mod ulus of the composite was determined from the linear portion of 3. Results and discussion the stress-strain curve Cyclic fiber push-in tests were performed using a desktop 3.1. Scanning Auger analysis apparatus previously described, but with the addition of a sym- metrically placed pair of capacitance gauges for displacement Elemental composition depth profiles obtained from scan- measurements. Thin sections of the composites, cut normal to ning Auger microprobe analysis for the BN/SiC coatings on the fiber axis with a diamond saw, and polished down to a 0. 1-um Hi-Nicalon fibers are shown in Fig. 1. The coating consists of finish on both top and bottom faces were tested. Final specimen 0. 15 um thick Si-rich SiC followed by 0.6 um of carbon rich thickness was typically about 3 mm. Fibers were pushed in using"BN. In addition, unintentionally deposited carbon layer is also a 700-included-angle conical diamond indenter with a 10-um present between the SiC and"BN"coatings. Another predom diameter flat base. To prevent the sides of the conical inden- nantly carbon layer is also seen between the"BN"coating and ter from impacting the matrix, push-in distances were restricted the fiber surface. Presence of free Si has also been detected to just a couple of microns. Unless otherwise noted, each test in the SiC coating layer by Raman microspectroscopy. This is consisted of five cycles of loading and unloading between a consistent with the results of another study22 which found the selected maximum load and a minimum load of 0.01 N at room Sic layer to be rich in Si from scanning Auger analysi mperature in ambient atmosphere Thermogravimetric analysis(TGA)was carried out at a heat- 3.2. Microstructural analy ing rate of 5C/min under flowing air from room temperature to faced with a computerized data acquisition and analysis system. SEM micrographs taken from the polished cross-section of 1500C using a PerkinElmer TGA-7 system, which was inter- he unidirectional hot pressed composite are shown in Fig. 2

N.P. Bansal / Journal of the European Ceramic Society 29 (2009) 525–535 527 Fig. 3. TGA curves for BSAS monolith, BN/SiC-coated Hi-Nicalon fiber and Hi-Nicalon/BN/SiC/BSAS composite recorded at a heating rate of 5 ◦C/min in air. exposures in air and mechanical testing. The volume fraction of fibers in the composite was found to be ∼0.32. For high-temperature annealing, the CMC bars were rested on the edges of an alumina boat placed inside a programmable box furnace. The furnace temperature was raised at a heating rate of 20 ◦C/min. CMC bars were annealed at 550, 800, 900, 1000, 1100, and 1200 ◦C for 100 h in stagnant ambient air and furnace cooled. Dimensions and weight of each test bar were recorded before and after annealing. Mechanical properties were determined from apparent stress–strain curves recorded from a three-point flexure test using an Instron 4505 universal testing machine at a cross-head speed of 1.27 mm/min (0.05 in./min) and support span (L) of 40 mm. Strain gauges were glued to the tensile surfaces of the flexure test bars. Stress, σ, was calculated from beam theory, assuming a linear elastic beam, using the equation: σ = 3PL 2bh2 (1) where b and h are the width and thickness of the test sample and P is the load. The yield stress, σy, was taken from the onset of deviation from linearity in the stress–strain curve. Elastic mod￾ulus of the composite was determined from the linear portion of the stress–strain curve. Cyclic fiber push-in tests were performed using a desktop apparatus previously described,21 but with the addition of a sym￾metrically placed pair of capacitance gauges for displacement measurements. Thin sections of the composites, cut normal to the fiber axis with a diamond saw, and polished down to a 0.1-m finish on both top and bottom faces were tested. Final specimen thickness was typically about 3 mm. Fibers were pushed in using a 70◦-included-angle conical diamond indenter with a 10-m diameter flat base. To prevent the sides of the conical inden￾ter from impacting the matrix, push-in distances were restricted to just a couple of microns. Unless otherwise noted, each test consisted of five cycles of loading and unloading between a selected maximum load and a minimum load of 0.01 N at room temperature in ambient atmosphere. Thermogravimetric analysis (TGA) was carried out at a heat￾ing rate of 5 ◦C/min under flowing air from room temperature to 1500 ◦C using a PerkinElmer TGA-7 system, which was inter￾faced with a computerized data acquisition and analysis system. X-ray diffraction (XRD) patterns were recorded at room tem￾perature using a step scan procedure (0.02◦/2θ step, time/step 0.5 or 1 s) on a Phillips ADP-3600 automated diffractometer equipped with a crystal monochromator employing Cu K radi￾ation. Density was measured from dimensions and mass as well as by the Archimedes method. Microstructures of the polished cross-sections and fracture surfaces were observed in a JEOL JSM-840A scanning electron microscope. Prior to analysis, a thin layer of carbon was evaporated onto the SEM specimens for electrical conductivity. The elemental compositions of the fiber surface coatings were analyzed with a scanning Auger microprobe (Fisons Instru￾ments Microlab Model 310-F). The fibers for this analysis were mounted on a stainless steel sample mount by tacking the ends with colloidal graphite. Depth profiling was per￾formed by sequential ion-beam sputtering and Auger analysis. The ion etching was done with 3 keV argon ions rastered over an approximately 1 mm2 area. The etch rate in Ta2O5 under these conditions was 0.05 nm/s. Auger electron spec￾troscopy (AES) analysis of the coated Hi-Nicalon fibers was performed using an electron beam current of approximately 1.5 nA. The beam was rastered over a 2 m × 20m area of the fiber with the long axis of the area aligned with the long fiber axis. Spectra were acquired in integral mode at beam energy of 2 keV and depth profiles were generated by plot￾ting elemental peak areas against ion etch time. The atomic concentrations were calculated by dividing the peak areas by the spectrometer transmission function and the sensitivity fac￾tors for each peak, then scaling the results to total 100%. The sensitivity factors were derived from spectra of ion etched Si, B, SiC, BN, and TiO2 standards. The depth scale is from the Ta2O5 calibration and no attempt has been made to adjust for the actual etch rate for each material. Only the fibers with a smooth surface coating, rather than those having thick and rough coating morphologies, were used for Auger analy￾sis. 3. Results and discussion 3.1. Scanning Auger analysis Elemental composition depth profiles obtained from scan￾ning Auger microprobe analysis for the BN/SiC coatings on Hi-Nicalon fibers are shown in Fig. 1. The coating consists of ∼0.15m thick Si-rich SiC followed by ∼0.6m of carbon rich “BN”. In addition, unintentionally deposited carbon layer is also present between the SiC and “BN” coatings. Another predomi￾nantly carbon layer is also seen between the “BN” coating and the fiber surface. Presence of free Si has also been detected13 in the SiC coating layer by Raman microspectroscopy. This is consistent with the results of another study22 which found the SiC layer to be rich in Si from scanning Auger analysis. 3.2. Microstructural analysis SEM micrographs taken from the polished cross-section of the unidirectional hot pressed composite are shown in Fig. 2.

N P Bansal / Journal of the European Ceramic Society 29(2009)525-535 No anneal 550°C 800°C 900°c 1000°C 1100° 1200° Fig 4. Optical photographs showing Hi-Nicalon/BNSiC/BSAS composite bars annealed for 100h in air at various temperatures Uniform fiber distribution and good matrix infiltration within 3.3. Thermogravimetric analysis the fiber tows are evident. Some matrix porosity is also present Some of the filaments are of irregular shape rather than having The TGA curve for the Hi-Nicalon/BN/SiC/BSAS composite circular cross-section. The manufacturer reports an average fiber with a fiber volume fraction of 0.43 is given in Fig. 3. Also diameter of -14 um, but a large variation in the diameter of the shown for comparison are the curves for the BN/SiC-coated Hi- filaments within a fiber tow can be seen. The BN/SiC surface Nicalon fibers and a BSAs monolithic sample hot pressed at oating has been detached from some of the fibers during metal- 1300C for 2 h at 4 ksi (27.6 MPa). The monolithic ceramic lography or composite processing Debonding or loss of the fiber exhibits hardly any weight change and appears to be stable up to oating may lead to adverse reactions between the fibers and the 1500 C in air. The composite shows a negligible weight change oxide matrix at high-temperature resulting in strong fiber-matrix up to 1150.C. The total weight gain at 1450C is also small (-0.3%) In contrast, the fibers initially loose 0.5% weight 100m) 100um Fig. 5. SEM micrographs of polished cross-sections of unidirectional Hi-Nicalon/BNSIC/BSAS composites annealed in air for 100 h at various temperatures: (a) as- fabricated,(b)l000°C,(c)1100°,and(d)1200°C

528 N.P. Bansal / Journal of the European Ceramic Society 29 (2009) 525–535 Fig. 4. Optical photographs showing Hi-Nicalon/BN/SiC/BSAS composite bars annealed for 100 h in air at various temperatures. Uniform fiber distribution and good matrix infiltration within the fiber tows are evident. Some matrix porosity is also present. Some of the filaments are of irregular shape rather than having circular cross-section. The manufacturer reports an average fiber diameter of ∼14m, but a large variation in the diameter of the filaments within a fiber tow can be seen. The BN/SiC surface coating has been detached from some of the fibers during metal￾lography or composite processing. Debonding or loss of the fiber coating may lead to adverse reactions between the fibers and the oxide matrix at high-temperature resulting in strong fiber–matrix bonding. 3.3. Thermogravimetric analysis The TGA curve for the Hi-Nicalon/BN/SiC/BSAS composite with a fiber volume fraction of 0.43 is given in Fig. 3. Also shown for comparison are the curves for the BN/SiC-coated Hi￾Nicalon fibers and a BSAS monolithic sample hot pressed at 1300 ◦C for 2 h at 4 ksi (∼27.6 MPa). The monolithic ceramic exhibits hardly any weight change and appears to be stable up to 1500 ◦C in air. The composite shows a negligible weight change up to ∼1150 ◦C. The total weight gain at 1450 ◦C is also small (∼0.3%). In contrast, the fibers initially loose ∼0.5% weight Fig. 5. SEM micrographs of polished cross-sections of unidirectional Hi-Nicalon/BN/SiC/BSAS composites annealed in air for 100 h at various temperatures: (a) as-fabricated, (b) 1000 ◦C, (c) 1100 ◦C, and (d) 1200 ◦C.

N P Bansal /Journal of the European Ceramic Sociery 29(2009)525-535 1mm|[ 100 uml(d) 100pm Fig. 6. SEM micrographs of polished cross-sections of unidirectional Hi-Nicalon/BN/SIC/BSAS composites annealed at 1200'C for 100 h in ai up to 850C, probably due to the loss of absorbed moisture. exposed surfaces. Pores were present in the surface layer. Signs This is followed by a large weight increase, possibly due to the of partial melting and gas bubble formation during heat treatment oxidation of BN into B2O3 and also SiC to SiO2, particularly at were also observed. From XRD analysis, both amorphous and higher temperatures. The total weight gain is found to be 3%0. celsian phases were detected in the surface layer. Since such a behavior was not observed23 in monolithic bSas material even 3. 4. Thermal ageing in air after heat treatment for 20 h in air at 1500oc. it is assumed to be caused by the presence of Hi-Nicalon fibers and the BN/SiC CMC bars were heat treated in ambient air for 100h at var- coating. In the presence of air, BN is probably oxidized to B2O3 lous temperatures. Optical photographs showing the physical which reacts with the matrix and/or silica formed from the oxi- appearance of the CMc bars before and after annealing at var 1000 ious temperatures are shown in Fig. 4. No changes in physical appearance were observed in samples annealed at 1000C or lower. However, the specimen annealed at 1100C was covered with a thin porous white layer that could be easily removed by polishing with a fine emery paper. The samples aged at 1200C were deformed and developed a thick shiny white layer on the MC+ HC(trace) 550°C Strain. 0, deg coated Hi-Nicalon fiber-reinforced celsian matrix composites annealed in air for Fig. 7. X-ray diffraction pattern taken from the surface of as hot-pressed. Hi- 100 h at various temperatures.( For interpretation of the references to color in Nicalon/BN/SiC/BSAS composite MC: monoclinic celsian: HC: hexacelsian. the artwork, the reader is referred to the web version of the article

N.P. Bansal / Journal of the European Ceramic Society 29 (2009) 525–535 529 Fig. 6. SEM micrographs of polished cross-sections of unidirectional Hi-Nicalon/BN/SiC/BSAS composites annealed at 1200 ◦C for 100 h in air. up to ∼850 ◦C, probably due to the loss of absorbed moisture. This is followed by a large weight increase, possibly due to the oxidation of BN into B2O3 and also SiC to SiO2, particularly at higher temperatures. The total weight gain is found to be ∼3%. 3.4. Thermal ageing in air CMC bars were heat treated in ambient air for 100 h at var￾ious temperatures. Optical photographs showing the physical appearance of the CMC bars before and after annealing at var￾ious temperatures are shown in Fig. 4. No changes in physical appearance were observed in samples annealed at 1000 ◦C or lower. However, the specimen annealed at 1100 ◦C was covered with a thin porous white layer that could be easily removed by polishing with a fine emery paper. The samples aged at 1200 ◦C were deformed and developed a thick shiny white layer on the Fig. 7. X-ray diffraction pattern taken from the surface of as hot-pressed. Hi￾Nicalon/BN/SiC/BSAS composite. MC: monoclinic celsian; HC: hexacelsian. exposed surfaces. Pores were present in the surface layer. Signs of partial melting and gas bubble formation during heat treatment were also observed. From XRD analysis, both amorphous and celsian phases were detected in the surface layer. Since such a behavior was not observed23 in monolithic BSAS material even after heat treatment for 20 h in air at 1500 ◦C, it is assumed to be caused by the presence of Hi-Nicalon fibers and the BN/SiC coating. In the presence of air, BN is probably oxidized to B2O3 which reacts with the matrix and/or silica formed from the oxi￾Fig. 8. Apparent stress–strain curves recorded in three-point flexure for BN/SiC￾coated Hi-Nicalon fiber-reinforced celsian matrix composites annealed in air for 100 h at various temperatures. (For interpretation of the references to color in the artwork, the reader is referred to the web version of the article.)

30 N P Bansal Journal of the European Ceramic Society 29(2009)525-535 Table I Mechanical properties of unidirectional Hi-Nicalon/BN/SiC/BSAS composite annealed at various temperatures for 100 h in air; Vt=0.32(#Hi-NIC-BSAS-6-24-97) Annealing temperature(C) Density(g/cm) Weight change after annealing Ec(GPa) MPa) (%) Ou(MPa) 0.091 50 14 171 69 3.04 0.092 19 14 Deformed 43% a Measured at room temperature in three-point flexo dation of silicon carbide fibers, resulting in low-melting glassy the top debonded layer where damaged broken pieces of the phase which migrates to the sample surface. The bubble for- fibers are also present. producing amorphous silica and volatile CO and CO? gased ation may be related to the oxidation of Hi-Nicalon fibe 3.5. X-ray diffraction analys Stability of BN in moisture and oxygen containing atmo- spheres is an intrinsic problem in the long-term use of this KRD pattern recorded from surface of the as-fabricated material as fiber-matrix interface. 4 At 700C, the sensitivity CMC panel is given in Fig. 7. It shows the presence of only to moisture is controlled by the crystalline structure. Also, for- monoclinic celsian along with a trace amount of hexacelsian. mation of boric acid is minimal below 800C. However, B2O3 Additional diffraction peaks at d values of 0.504 nm(20=17.5) reacts readily with water to form HBO2. Therefore, in the pres- and 0.312 nm(20=28.5%)were detected in samples annealed at ence of moisture, B2O3 will undergo significant weight loss. The 900, 1000, and 1100C. The peak at d=0.312 nm was much product of hydrolysis is predominantly metaboric acid(HBO2) stronger than the one at d=0.504 nm in the 900C annealed with traces of orthoboric acid(H3 BO3). In dry air, BN shows specimen where as the reverse was true for the Cmc annealed at minimal oxidation up to 800C. At higher temperatures, B2O3 1000 and 1100C. The phase corresponding to these peaks could glass is formed on its surface. B2O3 has a low vapor pres- not be identified. a white shiny and glassy layer was formed on sure(<2 x 10-3 Torr)and volatilizes slowly at temperatures less the surface of the sample annealed at 1200C. The surface layer than 1100C. Most of the Bn coatings on fibers are deposited was found to contain monoclinic celsian and an unidentified at relatively low temperatures(1000C) and generally are amorphous phase from XRD analysis contaminated with carbo n and oxvgen impurities. Also, these BN coatings consist of randomly oriented microcrystalline or 3.6. Mechanical properties turbostratic grains and lack well ordered microstructures. The stability of BN towards moisture and its resistance towards air Apparent stress-strain curves recorded in three-point flex oxidation depend on the degree of its crystallinity. + BN with ure for the unidirectional BSAS matrix composite reinforced large d(002) spacing is much more reactive towards moisture with BN/SiC-coated Hi-Nicalon fibers before and after thermal than those close to the hexagonal BN structure probably because aging in air at various temperatures to 1100C, are presented of its less densely packed basal planes implying weaker atomic in Fig. 8. In earlier studies,20 hot pressed BSAS monolithic bonding. The higher the value of d(002) interlayer spacing, material showed flexural strength of 130 MPa, elastic modulus the less crystalline the material. BN with an interlayer spacing of 96 GPa, and failed in a brittle mode, as expected. In contrast, of d(002)=0.335 nm(3.35 A), which is close to the theoreti- the composites show initial linear elastic behavior followed by cal value of 0.333 nm, showed significantly improved stability an extended region of load carrying capability beyond the ini- towards moisture and air tial deviation from linearity. This indicates load transfer to the The matrix layers on the surface of CMC specimens fibers beyond the proportional limit indicating graceful failure annealed at 1100 and 1200C appear to have cracked and and true composite behavior. Various room temperature mechan delaminated, respectively as seen in the SEM micrographs ical properties of the composites, before and after thermal aging (Fig. 5)taken from the polished cross-sections. A large dif- in air at various temperatures, are summarized in Table 1. The ference in the coefficients of thermal expansion (CTE) of measured elastic modulus of the CMC is in good agreement Hi-Nicalon fiber(3.5 x 10-6oC-)and the oxide matrix with a value of 150 GPa, calculated from the rule-of-mixtures (5. x 10-6oC-)may be responsible for the observed (Ec=VmEm +ViEf where V is the volume fraction and the sub- cracking and delamination. This would provide an easy path scripts c, m, and f refer to the composite, matrix, and fiber, for the ingression of oxygen to the fiber bundles and accelerate respectively) using Em=96 GPa- and Er= 270 GPa. ,There the degradation of fibers from oxidation. No such cracking or is no effect of thermal annealing in arup to 1"C on the values delamination was observed in samples annealed at lower tem- of elastic modulus, yield stress, yield strain, and ultimate stress peratures. Higher magnification SEM micrographs(Fig. 6)from of the composites Mechanical behavior of the CMC annealed the 1200C annealed specimens show the presence of gas bub- at 1200C could not be recorded as this specimen had badly bles on its surface. Severe damage is also observed underneath deformed

530 N.P. Bansal / Journal of the European Ceramic Society 29 (2009) 525–535 Table 1 Mechanical propertiesa of unidirectional Hi-Nicalon/BN/SiC/BSAS composite annealed at various temperatures for 100 h in air; Vf = 0.32 (#Hi-NIC-BSAS-6-24-97) Annealing temperature (◦C) Density (g/cm3) Weight change after annealing Ec (GPa) σy (MPa) εy (%) σu (MPa) – 3.09 ± 0.03 – 137 122 0.091 759 550 3.12 None 145 155 0.108 853 800 3.06 None 150 138 0.096 814 900 3.16 None 151 171 0.114 769 1000 3.04 None 146 134 0.092 819 1100 2.90 None 142 143 0.102 736 1200 Deformed +0.43% a Measured at room temperature in three-point flexure. dation of silicon carbide fibers, resulting in low-melting glassy phase which migrates to the sample surface. The bubble for￾mation may be related to the oxidation of Hi-Nicalon fibers producing amorphous silica and volatile CO and CO2 gases. Stability of BN in moisture and oxygen containing atmo￾spheres is an intrinsic problem in the long-term use of this material as fiber–matrix interface.24 At 700 ◦C, the sensitivity to moisture is controlled by the crystalline structure. Also, for￾mation of boric acid is minimal below 800 ◦C. However, B2O3 reacts readily with water to form HBO2. Therefore, in the pres￾ence of moisture, B2O3 will undergo significant weight loss. The product of hydrolysis is predominantly metaboric acid (HBO2) with traces of orthoboric acid (H3BO3). In dry air, BN shows minimal oxidation up to 800 ◦C. At higher temperatures, B2O3 glass is formed on its surface. B2O3 has a low vapor pres￾sure (<2 × 10−3 Torr) and volatilizes slowly at temperatures less than 1100 ◦C. Most of the BN coatings on fibers are deposited at relatively low temperatures (∼1000 ◦C) and generally are contaminated with carbon and oxygen impurities. Also, these BN coatings consist of randomly oriented microcrystalline or turbostratic grains and lack well ordered microstructures. The stability of BN towards moisture and its resistance towards air oxidation depend on the degree of its crystallinity.24 BN with large d (0 0 2) spacing is much more reactive towards moisture than those close to the hexagonal BN structure probably because of its less densely packed basal planes implying weaker atomic bonding. The higher the value of d (0 0 2) interlayer spacing, the less crystalline the material. BN with an interlayer spacing of d (0 0 2) = 0.335 nm (3.35 Å), which is close to the theoreti￾cal value of 0.333 nm, showed significantly improved stability towards moisture and air. The matrix layers on the surface of CMC specimens annealed at 1100 and 1200 ◦C appear to have cracked and delaminated, respectively as seen in the SEM micrographs (Fig. 5) taken from the polished cross-sections. A large dif￾ference in the coefficients of thermal expansion (CTE) of Hi-Nicalon fiber (∼3.5 × 10−6 ◦C−1) and the oxide matrix20 (∼5.28 × 10−6 ◦C−1) may be responsible for the observed cracking and delamination. This would provide an easy path for the ingression of oxygen to the fiber bundles and accelerate the degradation of fibers from oxidation. No such cracking or delamination was observed in samples annealed at lower tem￾peratures. Higher magnification SEM micrographs (Fig. 6) from the 1200 ◦C annealed specimens show the presence of gas bub￾bles on its surface. Severe damage is also observed underneath the top debonded layer where damaged broken pieces of the fibers are also present. 3.5. X-ray diffraction analysis XRD pattern recorded from surface of the as-fabricated CMC panel is given in Fig. 7. It shows the presence of only monoclinic celsian along with a trace amount of hexacelsian. Additional diffraction peaks at d values of 0.504 nm (2θ = 17.5◦) and 0.312 nm (2θ = 28.5◦) were detected in samples annealed at 900, 1000, and 1100 ◦C. The peak at d = 0.312 nm was much stronger than the one at d = 0.504 nm in the 900 ◦C annealed specimen where as the reverse was true for the CMC annealed at 1000 and 1100 ◦C. The phase corresponding to these peaks could not be identified. A white shiny and glassy layer was formed on the surface of the sample annealed at 1200 ◦C. The surface layer was found to contain monoclinic celsian and an unidentified amorphous phase from XRD analysis. 3.6. Mechanical properties Apparent stress–strain curves recorded in three-point flex￾ure for the unidirectional BSAS matrix composite reinforced with BN/SiC-coated Hi-Nicalon fibers, before and after thermal aging in air at various temperatures to 1100 ◦C, are presented in Fig. 8. In earlier studies11,20 hot pressed BSAS monolithic material showed flexural strength of 130 MPa, elastic modulus of 96 GPa, and failed in a brittle mode, as expected. In contrast, the composites show initial linear elastic behavior followed by an extended region of load carrying capability beyond the ini￾tial deviation from linearity. This indicates load transfer to the fibers beyond the proportional limit indicating graceful failure and true composite behavior. Various room temperature mechan￾ical properties of the composites, before and after thermal aging in air at various temperatures, are summarized in Table 1. The measured elastic modulus of the CMC is in good agreement with a value of 150 GPa, calculated from the rule-of-mixtures (Ec = VmEm + VfEf where V is the volume fraction and the sub￾scripts c, m, and f refer to the composite, matrix, and fiber, respectively) using Em = 96 GPa20 and Ef = 270 GPa.18,19 There is no effect of thermal annealing in air up to 1100 ◦C on the values of elastic modulus, yield stress, yield strain, and ultimate stress of the composites. Mechanical behavior of the CMC annealed at 1200 ◦C could not be recorded as this specimen had badly deformed

N P Bansal /Journal of the European Ceramic Sociery 29(2009)525-535 No Anneal 550°C 800°C 100um 100um 900°C 1000°c 1100°C 副-邮圆 100 Fig 9. SEM micrographs showing fracture surfaces of BNSiC-coated Hi-Nicalon fiber-reinforced celsian matrix composites annealed in air for 100 h at various 3.7. SEM of fracture surfaces toughening behavior for all the annealed CMC samples. Typical higher magnification SEM pictures showing fracture surfaces SEM micrographs of fracture surfaces of the composites, of the as-fabricated CMC and the one air-annealed at 1100C annealed for 100 h in air at various temperatures up to 1100C, are presented in Fig. 10. Debonding occurs primarily between fter the room temperature flexure test, are shown in Fig. 9. the fiber and the innermost coating in both specimens. In some Extensive long lengths of fiber pullout are observed indicating instances, the coating appears to stay with the fibers, particu Unannealed 1100° C Anneal 0 Fig. 10. SEM micrographs showing fracture surfaces of BN/SiC-coated Hi-Nicalon fiber-reinforced celsian matrix composites: as-fabricated and those annealed at

N.P. Bansal / Journal of the European Ceramic Society 29 (2009) 525–535 531 Fig. 9. SEM micrographs showing fracture surfaces of BN/SiC-coated Hi-Nicalon fiber-reinforced celsian matrix composites annealed in air for 100 h at various temperatures. 3.7. SEM of fracture surfaces SEM micrographs of fracture surfaces of the composites, annealed for 100 h in air at various temperatures up to 1100 ◦C, after the room temperature flexure test, are shown in Fig. 9. Extensive long lengths of fiber pullout are observed indicating toughening behavior for all the annealed CMC samples. Typical higher magnification SEM pictures showing fracture surfaces of the as-fabricated CMC and the one air-annealed at 1100 ◦C are presented in Fig. 10. Debonding occurs primarily between the fiber and the innermost coating in both specimens. In some instances, the coating appears to stay with the fibers, particu￾Fig. 10. SEM micrographs showing fracture surfaces of BN/SiC-coated Hi-Nicalon fiber-reinforced celsian matrix composites: as-fabricated and those annealed at 1100 ◦C for 100 h in air

N P Bansal Journal of the European Ceramic Society 29(2009)525-535 Table 2 =194 Effects of annealing on fiber push-in test resul directional Triction=10 MP Hi-Nicalon/BN/SIC/BSAS composite; 12 plies, Vi=0.32(#Hi-NIC-BSAS-6- 24-97) Annealing conditions Friction (MPa) Debond As-fab 0.31士0.14 100h, 0.53±0.47a 8.33±1.72a Measured for fibers in the interior of the annealed CMC sample Fiber displacement, 1 um/div Poisson expansion of the fibers leads to an overestimation of Load versus fiber displacement curve recorded during fiber T friction values, the relative changes in friction with load cycling as-fabricated celsian matrix composite reinforced with BNS could be followed using Eq (2). In addition, a debond initiation stress, od, could be calculated from the debond initiation load Fa (load at which fiber end begins to move during first loading larly in the unannealed sample. Surface of the debonded fibers cycle)by the relation appears to be smooth indicating no fiber degradation or chemi- cal reaction with the matrix during high-temperature annealing od atll00°C The results of fiber push-in data are summarized in Table 2. 3.8. Fiber-matrix interface The scatters in the values of debond stress (od) and frictional sliding stress(tf) are due to the accuracy of measurements In as-fabricated fiber-reinforced composites, the fibers may for a particular fiber as well from the variations in experience a several hundred MPa clamping force upon cool- ues from fiber-to-fiber. The variation in frictional sliding stress ing from the CMC processing temperature. Thermal mismatch (tr) during successive loading-unloading cycles was small and stresses result from smaller thermal expansion of reinforcing within the standard deviation. Values of debond stress, od, and fibers as compared to the matrix. A compliant layer is necessary frictional sliding stres to reduce the stresses. The importance of a compliant interface 0.31+0.14 GPa and 10.4+3. 1 MPa, respectively, compared yer for a strong and tough CMC has been emphasized by with 0.53+0.47 GPa and 8.33+1.72 MPa for the fibers in the various workers. -Both graphitic and pyrolytic carbon and interior of the 1200C annealed sample. These results indicate hexagonal or turbostratic BN have exceptionally low moduli. that only the outer ply of the 1200 C annealed CMC specimens Ceramic composites that demonstrate good damage tolerance has been degraded from oxidation whereas the bulk interior part generally contain carbon or BN layer between the fiber and remains unaffected matrix with some exceptions such as porous oxide matrix CMCs For strong and particularly tough CMCs, the fiber-matrix 3.9. Degradation mechanism at 1200oC interface must be sufficiently weak to allow debonding at the interface, yet strong enough for effective load transfer from the a possible mechanism explaining the various steps involved matrix to the fiber. Fiber-matrix debonding and frictional slid the degradation of Hi-Nicalon/BN/SiC/BSAS CMC on ing stresses at the fiber-matrix interface were evaluated from annealing in air at 1200Cispresented in Fig 12. During anneal fiber push-in tests. About 12-15 fibers/coupon were individu- ing at 1200 C, the surface matrix layer delaminates from the ally pushed in for the air annealed and as-fabricated composites. composite ply underneath(SEM micrograph of Fig. 5), probably A typical cyclic push-in curve at room temperature for the as- due to the large CTE mismatch between the Hi-Nicalon fibers cycle, is given in Fig. 11. The data were analyzed by first sub- gen into the composite. This causes the oxidation of SiC ibig eceived composite, along with reloading part of the second and the celsian matrix. This facilitates the ingression of ox tracting the appropriate load-train compliance correction from particularly those which have lost the duplex CVD coating(see the measured displacements. An estimate of frictional slidi Fig. 2), resulting in the formation of Sioz stress, friction, was obtained using the constant Friction model of Marshall and Oliver25 which includes effects of residual SiC()+O2(g)- SiO2(s)+Co(g)+ CO2(g) stresses, but does not consider fiber roughness or Poisson expan- sion. Value of friction was determined by fitting the compliance tion of surface matrix layer in FRC corrected data from each reloading curve to the relationship Oxidation of sic fibers O2→Si02+00+cO2 2r3 Reaction of SiOz and Celsian: Et friction Low m.p. ternary phase where u is the fiber end displacement, uo is the residual fiber end On cooling: formation of Celsian crystals in glass matrix displacement after the previous unloading, Fis the applied load, Fig. 12. Proposed mechanism for degradation of Hi-Nicalon/BN/SiC/BSAS rf is the fiberradius, and Ef is the fibermodulus. While neglecting composites during annealing at 1200C in ai

532 N.P. Bansal / Journal of the European Ceramic Society 29 (2009) 525–535 Fig. 11. Load versus fiber displacement curve recorded during fiber push-in test of as-fabricated celsian matrix composite reinforced with BN/SiC-coated Hi-Nicalon fibers. larly in the unannealed sample. Surface of the debonded fibers appears to be smooth indicating no fiber degradation or chemi￾cal reaction with the matrix during high-temperature annealing at 1100 ◦C. 3.8. Fiber–matrix interface In as-fabricated fiber-reinforced composites, the fibers may experience a several hundred MPa clamping force upon cool￾ing from the CMC processing temperature. Thermal mismatch stresses result from smaller thermal expansion of reinforcing fibers as compared to the matrix. A compliant layer is necessary to reduce the stresses. The importance of a compliant interface layer for a strong and tough CMC has been emphasized by various workers.26–29 Both graphitic and pyrolytic carbon and hexagonal or turbostratic BN have exceptionally low moduli. Ceramic composites that demonstrate good damage tolerance generally contain carbon or BN layer between the fiber and matrix with some exceptions such as porous oxide matrix CMCs. For strong and particularly tough CMCs, the fiber–matrix interface must be sufficiently weak to allow debonding at the interface, yet strong enough for effective load transfer from the matrix to the fiber. Fiber–matrix debonding and frictional slid￾ing stresses at the fiber–matrix interface were evaluated from fiber push-in tests. About 12–15 fibers/coupon were individu￾ally pushed in for the air annealed and as-fabricated composites. A typical cyclic push-in curve at room temperature for the as￾received composite, along with reloading part of the second cycle, is given in Fig. 11. The data were analyzed by first sub￾tracting the appropriate load-train compliance correction from the measured displacements. An estimate of frictional sliding stress, τfriction, was obtained using the constant τfriction model of Marshall and Oliver25 which includes effects of residual stresses, but does not consider fiber roughness or Poisson expan￾sion. Value of τfriction was determined by fitting the compliance corrected data from each reloading curve to the relationship u = u0 +  F2 8π2 r3 f Ef τfriction  (2) where u is the fiber end displacement, u0 is the residual fiber end displacement after the previous unloading, F is the applied load, rf is the fiber radius, andEf is the fiber modulus. While neglecting Table 2 Effects of annealing on fiber push-in test results for unidirectional Hi-Nicalon/BN/SiC/BSAS composite; 12 plies, Vf = 0.32 (#Hi-NIC-BSAS-6- 24-97) Annealing conditions σd (GPa) τfriction (MPa) As-fabricated 0.31 ± 0.14 10.4 ± 3.1 100 h, air, 1200 ◦C 0.53 ± 0.47a 8.33 ± 1.72a a Measured for fibers in the interior of the annealed CMC sample. Poisson expansion of the fibers leads to an overestimation of τfriction values, the relative changes in τfriction with load cycling could be followed using Eq. (2). In addition, a debond initiation stress, σd, could be calculated from the debond initiation load, Fd (load at which fiber end begins to move during first loading cycle) by the relation σd = Fd πr2 f (3) The results of fiber push-in data are summarized in Table 2. The scatters in the values of debond stress (σd) and frictional sliding stress (τf) are due to the accuracy of measurements for a particular fiber as well as from the variations in val￾ues from fiber-to-fiber. The variation in frictional sliding stress (τf) during successive loading–unloading cycles was small and within the standard deviation. Values of debond stress, σd, and frictional sliding stress, τf, for the as-fabricated CMC were 0.31 ± 0.14 GPa and 10.4 ± 3.1 MPa, respectively, compared with 0.53 ± 0.47 GPa and 8.33 ± 1.72 MPa for the fibers in the interior of the 1200 ◦C annealed sample. These results indicate that only the outer ply of the 1200 ◦C annealed CMC specimens has been degraded from oxidation whereas the bulk interior part remains unaffected. 3.9. Degradation mechanism at 1200 ◦C A possible mechanism explaining the various steps involved in the degradation of Hi-Nicalon/BN/SiC/BSAS CMC on annealing in air at 1200 ◦C is presented in Fig. 12. During anneal￾ing at 1200 ◦C, the surface matrix layer delaminates from the composite ply underneath (SEM micrograph of Fig. 5), probably due to the large CTE mismatch between the Hi-Nicalon fibers and the celsian matrix. This facilitates the ingression of oxy￾gen into the composite. This causes the oxidation of SiC fibers, particularly those which have lost the duplex CVD coating (see Fig. 2), resulting in the formation of SiO2: SiC (s) + O2 (g) → SiO2 (s) + CO (g) + CO2 (g) (4) Fig. 12. Proposed mechanism for degradation of Hi-Nicalon/BN/SiC/BSAS composites during annealing at 1200 ◦C in air

N P Bansal /Journal of the European Ceramic Sociery 29(2009)525-535 wt%Al203 Tridymite A2O3:15% 1400° Mullite /8 BS- B2s YY￥ Wt Al2O3 Fig 13. Phase diagram of Bao-Al2O3-SiO2 system showing the presence of low-melting ternary eutectics. The silica formed reacts with celsian resulting in the formation SEM micrograph of Fig. 6. On cooling, celsian crystals pre- of a low-melting phase. The phase diagram of Ba0-Al2O3-SiOz cipitate from the melt leaving behind a glassy matrix which ternary system is given in Fig. 13. It does show the pres- is richer in SiO2 and poorer in BaO and Al2O3 than celsian ence of a ternary phase with a melting point of 1122 C. This SEM micrograph(Fig. 14) taken from the surface of 1200C phase is richer in SiOz but poorer in BaO and Al2O3 than cel- annealed CMC shows the presence of elongated crystals in some sian. Formation of gaseous by-products CO and CO2 during glassy matrix. From qualitative EDS analysis, these crystals are reaction(4)results in the evolution of bubbles as observed in seen to be celsian and the glassy matrix is found to be richer SEM Picture a EDS Area 2 EDS Area 1 g surface of Hi-Nicalon/ BN/SiC/BSAS composite annealed at 1200 C for 100 h in air. Also shown are the EDS analyses of two areas. labeled as 2. of the CMC surface

N.P. Bansal / Journal of the European Ceramic Society 29 (2009) 525–535 533 Fig. 13. Phase diagram of BaO–Al2O3–SiO2 system showing the presence of low-melting ternary eutectics. The silica formed reacts with celsian resulting in the formation of a low-melting phase. The phase diagram of BaO–Al2O3–SiO2 ternary system30 is given in Fig. 13. It does show the pres￾ence of a ternary phase with a melting point of 1122 ◦C. This phase is richer in SiO2 but poorer in BaO and Al2O3 than cel￾sian. Formation of gaseous by-products CO and CO2 during reaction (4) results in the evolution of bubbles as observed in SEM micrograph of Fig. 6. On cooling, celsian crystals pre￾cipitate from the melt leaving behind a glassy matrix which is richer in SiO2 and poorer in BaO and Al2O3 than celsian. SEM micrograph (Fig. 14) taken from the surface of 1200 ◦C annealed CMC shows the presence of elongated crystals in some glassy matrix. From qualitative EDS analysis, these crystals are seen to be celsian and the glassy matrix is found to be richer Fig. 14. SEM micrograph showing surface of Hi-Nicalon/BN/SiC/BSAS composite annealed at 1200 ◦C for 100 h in air. Also shown are the EDS analyses of two different areas, labeled as 1 and 2, of the CMC surface

N P Bansal Journal of the European Ceramic Society 29(2009)525-535 MC+ HC(trace) Powder scrapped 广HC 20, deg Fig 15. X-ray diffraction spectra from surfaces of Hi-Nicalon/BN/SiC/BSAS composites:(a) as-fabricated and(b) annealed at 1200 C for 100h in air. in SiOz but poorer in Bao and Al2O3 than celsian. The XRD 5. Conclusions patterns taken from the as-fabricated CMC and that for 100h in air at 1200C are given in Fig. 15. Onl Celsian matrix composites reinforced with BN/SiC-coated phase is detected in the as-fabricated sample wherea Hi-Nicalon fibers are stable in air up to 1100C. However, their ind some amorphous phase are present in the 1200C-annealed mechanical properties are severely degraded at higher tempera CMC. Formation of a low-melting glass phase has also been tures due to oxidation er during the study of BsAs environmental ba rier coating(EBC)on Si-based ceramic substrates such as CV Acknowledgments SiC and SiCr/SiC composite. On heat treatment, the plasma sprayed BSAS coating reacted with the silica layer, formed from Thanks are due to John Setlock, Ron Phillips, Terry Kacik, oxidation of the Si-based ceramic substrate, resulting in low- and Ralph Garlick for their technical assistance during compos melting ternary glass phase which was found to be richer in ite processing and characterization. Don Wheeler assisted with SiOz but poorer in BaO and Al2O3 than celsian. Formation of the scanning Auger analysis and Jeff Eldridge with fiber push-in amorphous silica, due to the oxidation of SiC whiskers, has also testing. This research was supported by NASAs Ultra Efficient been reported during oxidation study of Sic whisker reinforced Engine Technology (UEET) Project and Hypersonic Project of mullite/zirconia composites at 1000-1350C. At 1200.c the Fundamental Aeronautics Progran or higher temperatures, formation of zircon was observed from the reaction between ZrO2 and Sio2. Secondary mul- References lite were also formed through a solution-reprecipita mechanism I. Bansal, N. P, ed, Handbook of Ceramic Composites. Kluwer Academic 2. Boccaccini, A. R, Continuous fiber reinforced glass and glass-ceramic matrix composites. In Handbook of Ceramic Composites, ed. N. P. Ban Kluwer Academic Publishers, Boston(MA), 2005, Pp. 461-484 <. Room temperature mechanical properties of BN/SiC-coated 3. Bansal, N.P., Ceramic fiber-reinforced glass-ceramic matrix composite. US Hi-Nicalon fiber-reinforced celsian matrix composites remained Patent 5,214004,May25,1993 unaffected after thermal aging for 100 h in air at various tem- 4. Bansal, N P, Method of producing a ceramic fiber-reinforced glass-ceramic peratures up to 1100oC. A thin white layer had formed on matrix composite. US Patent 5, 281, 559, January 25, 1994 the surface of the 1100C annealed sample and its density 5. Bansal. N. P. CVD SiC fiber-reinforced barium aluminosilicate glass-ceramic matrix composites. Mater Sci. Eng. A, 1996. 220(1-2), decreased from 3.09 to 2.90g/cm. However, the specimen 129-139. annealed at 1200C gained 0.43% weight, deformed in shape 6. Bansal, N.P., McCluskey, P, Linsey, G, Murphy, Dand Levan,G, Nicalon and size, and was covered with a thick shiny white porous layer fiber-reinforced celsian glass-ceramic matrix composites In Proceedin that could be easily peeled off. From X-ray diffraction analy f Annual HITEMP Review, Vol. Ill, 1995, NASA CP 10178, P. 41 sis, this surface layer was found to consist of amorphous and monoclinic celsian phases. The fibers in this surface layer had 7. Bansal, N. P, SiC fiber-reinforced celsian composites. In Handbook of Ceramic Composites, ed N. P. Bansal. Kluwer Academic Publishers, Boston broken into small pieces. The fiber-matrix interface in the inte- (MA),2005,pp.227-249 rior of the coupons was characterized through fiber push-in 8. Bansal, N P and Drummond Ill,CH,Kinetics of hexacelsian-to-celsian hnique. Values of debond stress, od, and frictional sliding phase transformation in SrAl2Si2O8. J. Am. Ceram. Soc., 1993, 76(5), stress, tf, for the as-fabricated CMC were 0.31+0.14 GPa and 9. Bansal. N. P. Mechanical behavior of silicon carbide fiber-reinforced stron- 0.4+3.1 MPa, respectively, compared with 0.53+0.47 GPa and 8. 33+ 1.72 MPa for the fibers in the interior of the 1200oC annealed sample indicating hardly any change in fiber-matrix 10. Bansal, N P and k J. A, Fabrication of fiber-reinforced celsian matrix interface Microstructures of the annealed specimens were inves- omposites. Composites: Part A, 2001, 32, 1021-1029 tigated using SEM. Only the surface ply of the 1200 C annealed 11.Bansal, N. P, Strong and tough Hi-Nicalon fiber-reinforced celsian matrix omposites. J. Am. Ceram. Soc., 1997, 80(9), 2407-2409. specimens had degraded from oxidation whereas the bulk inte- 12. Gyekenyesi, J.Z. and Bansal, N P. High temperature mechanical properties rior part of the CMC was unaffected of Hi-Nicalon fiber-reinforced celsian composites In Advances in Ceramic

534 N.P. Bansal / Journal of the European Ceramic Society 29 (2009) 525–535 Fig. 15. X-ray diffraction spectra from surfaces of Hi-Nicalon/BN/SiC/BSAS composites: (a) as-fabricated and (b) annealed at 1200 ◦C for 100 h in air. in SiO2 but poorer in BaO and Al2O3 than celsian. The XRD patterns taken from the as-fabricated CMC and that annealed for 100 h in air at 1200 ◦C are given in Fig. 15. Only celsian phase is detected in the as-fabricated sample whereas celsian and some amorphous phase are present in the 1200 ◦C-annealed CMC. Formation of a low-melting glass phase has also been reported31 earlier during the study of BSAS environmental bar￾rier coating (EBC) on Si-based ceramic substrates such as CVD SiC and SiCf/SiC composite. On heat treatment, the plasma sprayed BSAS coating reacted with the silica layer, formed from oxidation of the Si-based ceramic substrate, resulting in low￾melting ternary glass phase which was found to be richer in SiO2 but poorer in BaO and Al2O3 than celsian. Formation of amorphous silica, due to the oxidation of SiC whiskers, has also been reported during oxidation study of SiC whisker reinforced mullite/zirconia composites32 at 1000–1350 ◦C. At 1200 ◦C or higher temperatures, formation of zircon was observed from the reaction between ZrO2 and SiO2. Secondary mul￾lite grains were also formed through a solution-reprecipitation mechanism. 4. Summary Room temperature mechanical properties of BN/SiC-coated Hi-Nicalon fiber-reinforced celsian matrix composites remained unaffected after thermal aging for 100 h in air at various tem￾peratures up to 1100 ◦C. A thin white layer had formed on the surface of the 1100 ◦C annealed sample and its density decreased from 3.09 to 2.90 g/cm3. However, the specimen annealed at 1200 ◦C gained 0.43% weight, deformed in shape and size, and was covered with a thick shiny white porous layer that could be easily peeled off. From X-ray diffraction analy￾sis, this surface layer was found to consist of amorphous and monoclinic celsian phases. The fibers in this surface layer had broken into small pieces. The fiber–matrix interface in the inte￾rior of the coupons was characterized through fiber push-in technique. Values of debond stress, σd, and frictional sliding stress, τf, for the as-fabricated CMC were 0.31 ± 0.14 GPa and 10.4 ± 3.1 MPa, respectively, compared with 0.53 ± 0.47 GPa and 8.33 ± 1.72 MPa for the fibers in the interior of the 1200 ◦C annealed sample indicating hardly any change in fiber–matrix interface. Microstructures of the annealed specimens were inves￾tigated using SEM. Only the surface ply of the 1200 ◦C annealed specimens had degraded from oxidation whereas the bulk inte￾rior part of the CMC was unaffected. 5. Conclusions Celsian matrix composites reinforced with BN/SiC-coated Hi-Nicalon fibers are stable in air up to 1100 ◦C. However, their mechanical properties are severely degraded at higher tempera￾tures due to oxidation. Acknowledgments Thanks are due to John Setlock, Ron Phillips, Terry Kacik, and Ralph Garlick for their technical assistance during compos￾ite processing and characterization. Don Wheeler assisted with the scanning Auger analysis and Jeff Eldridge with fiber push-in testing. This research was supported by NASA’s Ultra Efficient Engine Technology (UEET) Project and Hypersonic Project of the Fundamental Aeronautics Program. References 1. Bansal, N. P., ed., Handbook of Ceramic Composites. Kluwer Academic Publishers, Boston (MA), 2005. 2. Boccaccini, A. R., Continuous fiber reinforced glass and glass–ceramic matrix composites. In Handbook of Ceramic Composites, ed. N. P. Bansal. Kluwer Academic Publishers, Boston (MA), 2005, pp. 461–484. 3. Bansal, N. P., Ceramic fiber-reinforced glass–ceramic matrix composite. US Patent 5,214,004, May 25, 1993. 4. Bansal, N. P., Method of producing a ceramic fiber-reinforced glass–ceramic matrix composite. US Patent 5,281,559, January 25, 1994. 5. Bansal, N. P., CVD SiC fiber-reinforced barium aluminosilicate glass–ceramic matrix composites. Mater. Sci. Eng. A, 1996, 220(1–2), 129–139. 6. Bansal, N. P., McCluskey, P., Linsey, G., Murphy, D. and Levan, G., Nicalon fiber-reinforced celsian glass–ceramic matrix composites. In Proceedings of Annual HITEMP Review, Vol. III, 1995, NASA CP 10178, p. 41-1– 14. 7. Bansal, N. P., SiC fiber-reinforced celsian composites. In Handbook of Ceramic Composites, ed. N. P. Bansal. Kluwer Academic Publishers, Boston (MA), 2005, pp. 227–249. 8. Bansal, N. P. and Drummond III, C. H., Kinetics of hexacelsian-to-celsian phase transformation in SrAl2Si2O8. J. Am. Ceram. Soc., 1993, 76(5), 1321–1324. 9. Bansal, N. P., Mechanical behavior of silicon carbide fiber-reinforced stron￾tium aluminosilicate glass–ceramic composites. Mater. Sci. Eng. A, 1997, 231(1–2), 117–127. 10. Bansal, N. P. and Setlock, J. A., Fabrication of fiber-reinforced celsian matrix composites. Composites: Part A, 2001, 32, 1021–1029. 11. Bansal, N. P., Strong and tough Hi-Nicalon fiber-reinforced celsian matrix composites. J. Am. Ceram. Soc., 1997, 80(9), 2407–2409. 12. Gyekenyesi, J. Z. and Bansal, N. P. High temperature mechanical properties of Hi-Nicalon fiber-reinforced celsian composites. In Advances in Ceramic