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 composite 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 HiNicalon fiber-reinforced celsian matrix composites have been reported earlier.6–17 During high-temperature use, CMC components 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 celsian 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 properties 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 pushin 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 HiNicalon 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 precursor 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 nominal 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) composition 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 substitution 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 fabrication 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” composite. 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 modulus 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 symmetrically 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 indenter 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 heating rate of 5 ◦C/min under flowing air from room temperature to 1500 ◦C using a PerkinElmer TGA-7 system, which was interfaced with a computerized data acquisition and analysis system. X-ray diffraction (XRD) patterns were recorded at room temperature 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 radiation. 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 Instruments 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 performed 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 spectroscopy (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 plotting 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 factors 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 analysis. 3. Results and discussion 3.1. Scanning Auger analysis Elemental composition depth profiles obtained from scanning 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 predominantly 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 metallography 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 HiNicalon 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 various temperatures. Optical photographs showing the physical appearance of the CMC bars before and after annealing at various 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. HiNicalon/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 oxiFig. 8. Apparent stress–strain curves recorded in three-point flexure for BN/SiCcoated 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.)