《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_Thermal-shock behavior of a Nicalon-fiber-reinforced hybrid glass-ceramic composite.pdf_大学文库

COMPOSITES SCIENCE AND TECHNOLOGY ELSEⅤIER Composites Science and Technology 61(2001)1923-1930 www.elsevier.com/locate/compscitech Thermal-shock behavior of a Nicalon-fiber-reinforced hybrid glass-ceramic composite N. Chawla.*, K.K. Chawla, M. Koopman, B Patel, C. Coffin, J.I. Eldridge a Department of Chemical and Materials Engineering, Arizona State University, PO Box 876006, Tempe, AZ 85287-6006, USA Department of Materials and Mechanical Engineering, University of Alabama at Birmingham BEC 254, 1530 3rd Av. South, birmingham, AL 35294, US.A "NASA Glenn Research Center, MS 106-5, 21000 Brookpark Road, Cleveland, OH 44135, USA Received 1l January 2001; received sed form 5 June 2001; accepted 3 July 2001 Abstract A Nicalon-fiber-reinforced hybrid composite with a matrix of barium magnesium aluminosilicate(BMAs) glass with silicon carbide whiskers was subjected to thermal shock from elevated to ambient temperatures. The combination of Sic whisker and BMAS glass resulted in a hybrid matrix with a lower thermal expansion than that of the fibers, inducing tensile stresses in the fiber upon thermal shock. This stress state resulted in microstructural damage in the form of fiber cracking and cracking along the fiber/ matrix interface, as opposed to the conventional matrix cracking which is typically observed in ceramic-matrix composites. Sig- nificant damage in the composite was only observed after three thermal shock cycles. Flexural resonance measurements, used to evaluate thermal shock-induced changes in Youngs modulus, showed a reduction in modulus that correlated well with the onset of microstructural damage. Finally, fiber push-out tests, performed to evaluate changes in fiber/matrix interface strength after thermal cycling, indicated a slight decrease in interfacial strength, which was attributed to recession of the carbon-rich fiber surface during hermal shock. C 2001 Elsevier Science Ltd. All rights reserved Keywords: A Ceramic matrix composites; Whisker; Thermal shock; Damage; Fiber cracking 1. Introduction role in determining the toughness of the composite [1, 3- 7. While several CMC systems have very high strength Fiber-reinforced glass and glass-ceramic composites on account of load transfer and toughening mechanisms constitute a class of materials suitable for applications provided by a tailored fiber/matrix interface, many sys requiring a combination of lightweight, strength, and tems have a very low matrix cracking stress. This is not toughness at intermediate to elevated temperatures [1]. a desirable attribute since embrittlement in aggressive The glass or glass-ceramic matrix provides low density, environments may take place as a result of oxidation at while the fibers contribute to strength, stiffness, and the interface by oxygen ingression through the cracks toughness. Following the onset of matrix cracking in the Furthermore the higher the matrix crack stress, the composite, the presence of a weak fiber /matrix interface higher the allowable design stress for a given compo- at the tip of the growing crack leads to toughening nent. The incorporation of whiskers in the glass matrix through mechanisms such as crack blunting and deflec. can significantly increase the stress required for matrix tion, which are crucial in providing non-catastrophic cracking, and the fibers, with appropriate interface tai failure in the composite [1, 2]. Other fiber/matrix inter- loring, still provide high strength, work of fracture, and face properties, such as chemical composition, presence non-catastrophic failure [8,91 of voids, microcracking, and the microstructural stress Under thermal shock conditions the mismatch in the state near the interface, have also been shown to play a coefficient of thermal expansion of the matrix and fibers can contribute significantly to microcracking and indir Corresponding author. Tel: + 1-480-965-24 ectly to degradation in fiber strength through oxidation Consequently, the thermal history of the material, as ell as the ptibility of the material to a thermally 0266-3538/01/ S.see front matter C 2001 Elsevier Science Ltd. All rights reserved. PII:S0266-3538(01)00096

Thermal-shock behavior of a Nicalon-fiber-reinforced hybrid glass-ceramic composite N. Chawlaa,*, K.K. Chawlab, M. Koopmanb, B. Patelb, C. Coffinb, J.I. Eldridgec a Department of Chemical and Materials Engineering, Arizona State University, PO Box 876006, Tempe, AZ 85287-6006, USA bDepartment of Materials and Mechanical Engineering, University of Alabama at Birmingham, BEC 254, 1530 3rd Av. South, Birmingham, AL 35294, USA c NASA Glenn Research Center, MS 106-5, 21000 Brookpark Road, Cleveland, OH 44135, USA Received 11 January 2001; received in revised form 5 June 2001; accepted 3 July 2001 Abstract A Nicalon-fiber-reinforced hybrid composite with a matrix of barium magnesium aluminosilicate (BMAS) glass with silicon carbide whiskers was subjected to thermal shock from elevated to ambient temperatures. The combination of SiC whisker and BMAS glass resulted in a hybrid matrix with a lower thermal expansion than that of the fibers, inducing tensile stresses in the fiber upon thermal shock. This stress state resulted in microstructural damage in the form of fiber cracking and cracking along the fiber/ matrix interface, as opposed to the conventional matrix cracking which is typically observed in ceramic-matrix composites. Significant damage in the composite was only observed after three thermal shock cycles. Flexural resonance measurements, used to evaluate thermal shock-induced changes in Young’s modulus, showed a reduction in modulus that correlated well with the onset of microstructural damage. Finally, fiber push-out tests, performed to evaluate changes in fiber/matrix interface strength after thermal cycling, indicated a slight decrease in interfacial strength, which was attributed to recession of the carbon-rich fiber surface during thermal shock. # 2001 Elsevier Science Ltd. All rights reserved. Keywords: A. Ceramic matrix composites; Whisker; Thermal shock; Damage; Fiber cracking 1. Introduction Fiber-reinforced glass and glass-ceramic composites constitute a class of materials suitable for applications requiring a combination of lightweight, strength, and toughness at intermediate to elevated temperatures [1]. The glass or glass-ceramic matrix provides low density, while the fibers contribute to strength, stiffness, and toughness. Following the onset of matrix cracking in the composite, the presence of a weak fiber/matrix interface at the tip of the growing crack leads to toughening through mechanisms such as crack blunting and deflection, which are crucial in providing non-catastrophic failure in the composite [1,2]. Other fiber/matrix interface properties, such as chemical composition, presence of voids, microcracking, and the microstructural stress state near the interface, have also been shown to play a role in determining the toughness of the composite [1,3– 7]. While several CMC systems have very high strength, on account of load transfer and toughening mechanisms provided by a tailored fiber/matrix interface, many systems have a very low matrix cracking stress. This is not a desirable attribute since embrittlement in aggressive environments may take place as a result of oxidation at the interface by oxygen ingression through the cracks. Furthermore, the higher the matrix crack stress, the higher the allowable design stress for a given component. The incorporation of whiskers in the glass matrix can significantly increase the stress required for matrix cracking, and the fibers, with appropriate interface tailoring, still provide high strength, work of fracture, and non-catastrophic failure [8,9]. Under thermal shock conditions, the mismatch in the coefficient of thermal expansion of the matrix and fibers can contribute significantly to microcracking and indirectly to degradation in fiber strength through oxidation. Consequently, the thermal history of the material, as well as the susceptibility of the material to a thermally 0266-3538/01/$ - see front matter # 2001 Elsevier Science Ltd. All rights reserved. PII: S0266-3538(01)00096-3 Composites Science and Technology 61 (2001) 1923–1930 www.elsevier.com/locate/compscitech * Corresponding author. Tel.: +1-480-965-2402; fax: +1-480-965- 0037. E-mail address: nchawla@asu.edu (N. Chawla)

1924 N. Chawla et al. Composites Science and Technology 61(2001)1923-1930 ynamic environment, play a pivotal role in determining where m is the mass of the specimen, ff is the resonant mechanical properties of the material. Thermal shock frequency of the bar under flexural loading, and b, L, studies of conventional glass and glass-ceramic matrix and t are the width, length, and thickness of the bar, composites have included Nicalon"M fibers within various respectively. Ti can be written as: matrices, including calcium aluminosilicate(CAS), lithium aluminosilicate (LAS), magnesium aluminosilicate T1=1+6.585(1+0.0752v+081092 (MAS), and borosilicate glasses [2, 3, 10-16]. For a given fiber, differences in matrix chemistry among the various composites can yield variations in interfacial properties 8340(1+0.2023+2173y2)(/L) and corresponding effects on thermal and mechanical 1.00+6338(1+0.1408v+1.536v2)t/L) properties. The present study is aimed at examining the thermal shock behavior of a novel Nicalon -fiber -rein- forced barium magnesium aluminosilicate(bMas)glass- where v is the Poissons ratio of the material [18] ceramic matrix containing SiC whiskers. The effects of Fiber push-out tests were conducted on polished cross thermal shock on the evolution of microstructural sections of approximately 2 mm thickness. Push-out damage, modulus, and interfacial shear strength of the was conducted on a micromechanical testing system, composite were investigated using a conical diamond indenter with a 10 um diameter flat region on the bottom of the indenter. Details of the fiber push-out test and the testing apparatus used in 2. Experimental procedure these experiments are given elsewhere [19, 20 Laminated unidirectional composites were pi by passing the fibers through a glass slurry containing 3. Results and discussion Sic whiskers(Corning Inc, Corning, NY). The indivi dual lamina were stacked and hot-pressed in a nitrogen The microstructure of the BMAS composite showed it atmosphere between 1400 and 1500C. The volume to be fully dense(Fig. 1). A relatively homogeneous dis- fraction of fibers and whiskers was measured to be 0.42 tribution of Nicalon fibers was observed. The glass-cera and 0.22, respectively. Table I shows selected material mic matrix contained cordierite as the primary glass properties for the composite constituents [1, 17]. Sam- phase, and a uniform distribution of SiC whiskers. Energy ples were sectioned from the as-received plate to speci- dispersive spectroscopy(EDS)(Fig. 2), revealed a homo- mens approximately 20 mm in length, 6 mm wide, and 2 geneous distribution of Si in the fibers and whiskers. mm thick, and polished prior to thermal shock, to iso- and a uniform distribution of Al, Ba, and Mg in the late polishing-induced fiber damage. Samples for ther- matrix of the composite. Areas of residual glassy phase mal shock were brought to temperature and allowed to in the matrix, perhaps unreacted during the hot pressing stabilize for 15 min prior to quenching in water at 25C. process, were also observed The Youngs modulus of each sample was determined Damage induced by thermal shock, in the form of y an impulse resonance technique GrindoSonic). cracking, was only observed after three thermal shocks After measuring the flexural resonant frequency(kHz) There was no damage observed after one or two cycles, of the sample, the Youngs modulus(GPa) was calcu- except for isolated fiber cracking on the surface. This can lated by using the relationship be attributed to the high thermal-induced tensile stresses in the near-surface region after thermal shock [14, 21] E=0.9465 (1) After three cycles, however, cracking was not confined to the surface region alone, but was also observed in the specimen interior. The fact that more than one thermal shock is needed to induce damage in the volume of the material indicates that a few thermal shocks are required Table I to induce incipient damage in the fibers, which results in Selected material properties for composite constituent [1, 17] well developed cracking after three cycles. Fig. 3 shows a region prior to and after thermal shock from 850C to oung s Coefficient Poisson room temperature. The pulled out fibers at the surface modulu raction are an artifact of polishing. The cracking seems toinitiate in the fibers, primarily transverse to the fiber axis. No longitudinal cracking in the fibers was observed. Fiber BMAS 0.25 cracking seems to be followed by propagation into the 0.22 Nicalon 0.42 matrix or along the fiber/matrix interface, see Figs. 4 and 5. Blissett et al. [11] also documented two types of

dynamic environment, play a pivotal role in determining mechanical properties of the material. Thermal shock studies of conventional glass and glass–ceramic matrix composites have included NicalonTM fibers within various matrices, including calcium aluminosilicate (CAS), lithium aluminosilicate (LAS), magnesium aluminosilicate (MAS), and borosilicate glasses [2,3,10–16]. For a given fiber, differences in matrix chemistry among the various composites can yield variations in interfacial properties and corresponding effects on thermal and mechanical properties. The present study is aimed at examining the thermal shock behavior of a novel Nicalon-fiber-reinforced barium magnesium aluminosilicate (BMAS) glass– ceramic matrix containing SiC whiskers. The effects of thermal shock on the evolution of microstructural damage, modulus, and interfacial shear strength of the composite were investigated. 2. Experimental procedure Laminated unidirectional composites were produced by passing the fibers through a glass slurry containing SiC whiskers (Corning Inc., Corning, NY). The individual lamina were stacked and hot-pressed in a nitrogen atmosphere between 1400 and 1500 C. The volume fraction of fibers and whiskers was measured to be 0.42 and 0.22, respectively. Table 1 shows selected material properties for the composite constituents [1,17]. Samples were sectioned from the as-received plate to specimens approximately 20 mm in length, 6 mm wide, and 2 mm thick, and polished prior to thermal shock, to isolate polishing-induced fiber damage. Samples for thermal shock were brought to temperature and allowed to stabilize for 15 min prior to quenching in water at 25 C. The Young’s modulus of each sample was determined by an impulse resonance technique (GrindoSonic). After measuring the flexural resonant frequency (kHz) of the sample, the Young’s modulus (GPa) was calculated by using the relationship: E ¼ 0:9465 mff 2 b L3 t3 T1 ð1Þ where m is the mass of the specimen, ff is the resonant frequency of the bar under flexural loading, and b, L, and t are the width, length, and thickness of the bar, respectively. T1 can be written as: T1 ¼ 1 þ 6:585 1 þ 0:0752 þ 081092 t L 4 8:340 1 þ 0:2023 þ 2:1732 ð Þ t=L 4 1:00 þ 6:338 1 þ 0:1408 þ 1:5362 ð Þð Þ t=L 2 " # ð2Þ where is the Poisson’s ratio of the material [18]. Fiber push-out tests were conducted on polished cross sections of approximately 2 mm thickness. Push-out was conducted on a micromechanical testing system, using a conical diamond indenter with a 10 mm diameter flat region on the bottom of the indenter. Details of the fiber push-out test and the testing apparatus used in these experiments are given elsewhere [19,20]. 3. Results and discussion The microstructure of the BMAS composite showed it to be fully dense (Fig. 1). A relatively homogeneous distribution of Nicalon fibers was observed. The glass–ceramic matrix contained cordierite as the primary glass phase, and a uniform distribution of SiC whiskers. Energy dispersive spectroscopy (EDS) (Fig. 2), revealed a homogeneous distribution of Si in the fibers and whiskers, and a uniform distribution of Al, Ba, and Mg in the matrix of the composite. Areas of residual glassy phase in the matrix, perhaps unreacted during the hot pressing process, were also observed. Damage induced by thermal shock, in the form of cracking, was only observed after three thermal shocks. There was no damage observed after one or two cycles, except for isolated fiber cracking on the surface. This can be attributed to the high thermal-induced tensile stresses in the near-surface region after thermal shock [14,21]. After three cycles, however, cracking was not confined to the surface region alone, but was also observed in the specimen interior. The fact that more than one thermal shock is needed to induce damage in the volume of the material indicates that a few thermal shocks are required to induce incipient damage in the fibers, which results in well developed cracking after three cycles. Fig. 3 shows a region prior to and after thermal shock from 850 C to room temperature. The pulled out fibers at the surface are an artifact of polishing. The cracking seems to initiate in the fibers, primarily transverse to the fiber axis. No longitudinal cracking in the fibers was observed. Fiber cracking seems to be followed by propagation into the matrix or along the fiber/matrix interface, see Figs. 4 and 5. Blissett et al. [11] also documented two types of Table 1 Selected material properties for composite constituent [1,17] Young’s modulus (GPa) Coefficient of thermal expansion (106 / C) Poisson’s ratio Volume fraction BMAS 120 2.5 0.25 0.36 SiC whisker 400 3.0 0.30 0.22 Nicalon 192 4.0 0.30 0.42 1924 N. Chawla et al. / Composites Science and Technology 61 (2001) 1923–1930

N Chawla et al. / Composites Science and Technology 61(2001)1923-1930 composite, since the whiskers are randomly distributed The use of the Sic whiskers in the low expansion The thermal expansion of the composite hybrid matrix, BMAS matrix also results in a composite with reduced Chm, is given by [22] thermal expansion, vis a vis that of the fiber. Using the upper and lower bounds for elastic modulus of the hybrid awVwKw+aBMAS VBMAS KBMAS (3) matrix, upper and lower bounds of the thermal expansion WkW+VBMAS KBMAS of the composite in the longitudinal direction were calcu- lated, using the following expression by Schapery [24] where K 5 and aw, Vw, aBMAs, and vBmas, are he thermal expansion and volume fraction of the whis- o m Em/m +arErkr ker and BMAS matrix, respectively. The elastic mod Emim+ erve ulus of the hybrid matrix was then calculated using the approach by Hashin and Shtrikman [23], that prescribes The difference in upper and lower bound of the elastic upper and lower bounds for the elastic moduli of the modulus of the hybrid matrix did not significantly affect composite. The upper and lower bounds on the bulk the thermal expansion of the composite, which was calcu modulus of the hybrid matrix, Khm, are given by [23] lated as 3.35x10-b/C. This compares very well with the experimentally determined value of 3. 0x10-/C [25] The fiber damage was evaluated in terms of a decrease upper =Kw+ (4) in Young's modulus of the composite In the 850C sam- ple after three cycles, a measured decrease in Youngs KBMAS - Kw 3Kw+4G modulus from 155 to 85 GPa was observed. Thermal shock after one cycle resulted in no decrease in modulus Khm. lower KBMAS which once again, correlated well with microstructural observations. In thermal cycling of this composite, the microstructures were essentially crack-free[26]. The non- destructive nature of modulus measurement would seem to be more attractive than destructive testing, e.g., flex- Kw-kbmas 3KBMAS + 4GBMAS ural testing after thermal shock [27]. It is important to (5) note that flexural testing may not adequately quantify damage in thermal shocked materials. because of the where G is the shear modulus and is given by G=x(+v. combination of tensile, compressive, and shear stresses The computed Youngs modulus of the hybrid matrix involved. Thus, while a predominantly tensile mode and the coefficient of thermal expansion, in comparison may operate in the undamaged composite, the presence to that of Nicalon fiber, are given in Table 2 of microcracks in the damaged state may reduce the In a conventional glass-ceramic matrix (with no interlaminar shear strength and the composite may fail whiskers) reinforced with SiC fibers, am will be higher in a shear mode [14]. Furthermore, shifting of the neu than af, so under thermal shock conditions the matrix tral axis in continuous fiber ceramic composites results would be expected to be in tension, and cracking due to in a significant enhancement of the strength [28, 29 thermal stresses should initiate in the matrix, Table 3 It is interesting to note that microstructural damage indicates that despite the addition of Sic whiskers to the and decrease in modulus are also very much affected by matrix(the Sic whiskers have a similar CtE to that of the hold time prior to thermal shock. Blissett et al. [10] BMAS matrix, shown in Table 3), the hybrid matrix still determined the thermal shock resistance of unidirec has a lower Cte than that of the fiber, so on quenching, tional and cross-plied Nicalon fiber reinforced CAs the fiber is in tension, and cracks initiate in the fiber. It is glass-ceramic composites. The composite specimens were important to point out that, from a design standpoint, allowed to age for one hour prior to shock. Their results loading of the fibers in tension, rather than the matrix, indicated a strength decrease at intermediate tempera would seem to be a more desirable loading configuration tures followed by a perceived"recovery"in strength because of the higher fiber strength Table 3 Table 2 Debond and sliding stresses of Nicalon fibers in bmas. as measured Elastic and thermal properties of hybrid matrix and Nicalon fiber by fiber pushout testing Youngs Coefficient of Volume fraction Fiber diameter nodulus thermal expansion (GPa)(10-6C) AS- received 179±29 16.4士5.1 124±4.6 Hybrid matrix Shocked,500°C 23.2±19 8.7±1.9 Nicalon fiber Shocked,950°C 214±24 13.2±6.8 9.1士3.9

composite, since the whiskers are randomly distributed. The thermal expansion of the composite hybrid matrix, hm, is given by [22]: hm ¼ wVwKw þ BMASVBMASKBMAS VwKw þ VBMASKBMAS ð3Þ where K ¼ E 3 1ð Þ 2 , and w, Vw, BMAS, and VBMAS, are the thermal expansion and volume fraction of the whisker and BMAS matrix, respectively. The elastic modulus of the hybrid matrix was then calculated using the approach by Hashin and Shtrikman [23], that prescribes upper and lower bounds for the elastic moduli of the composite. The upper and lower bounds on the bulk modulus of the hybrid matrix, Khm, are given by [23]: Khm; upper ¼ Kw þ 1 Vw 1 KBMAS Kw þ 3Vw 3Kw þ 4Gw 2 6 6 4 3 7 7 5 ð4Þ Khm; lower ¼ KBMAS þ Vw 1 Kw KBMAS þ 3 1ð Þ Vw 3KBMAS þ 4GBMAS 2 6 6 4 3 7 7 5 ð5Þ where G is the shear modulus and is given by G ¼ E 2 1ð Þ þ . The computed Young’s modulus of the hybrid matrix and the coefficient of thermal expansion, in comparison to that of Nicalon fiber, are given in Table 2. In a conventional glass–ceramic matrix (with no whiskers) reinforced with SiC fibers, m will be higher than f, so under thermal shock conditions the matrix would be expected to be in tension, and cracking due to thermal stresses should initiate in the matrix. Table 3 indicates that despite the addition of SiC whiskers to the matrix (the SiC whiskers have a similar CTE to that of the BMAS matrix, shown in Table 3), the hybrid matrix still has a lower CTE than that of the fiber, so on quenching, the fiber is in tension, and cracks initiate in the fiber. It is important to point out that, from a design standpoint, loading of the fibers in tension, rather than the matrix, would seem to be a more desirable loading configuration because of the higher fiber strength. The use of the SiC whiskers in the low expansion BMAS matrix also results in a composite with reduced thermal expansion, vis a` vis that of the fiber. Using the upper and lower bounds for elastic modulus of the hybrid matrix, upper and lower bounds of the thermal expansion of the composite in the longitudinal direction were calculated, using the following expression by Schapery [24]: cl ¼ mEmVm þ fEfVf EmVm þ EfVf ð6Þ The difference in upper and lower bound of the elastic modulus of the hybrid matrix did not significantly affect the thermal expansion of the composite, which was calculated as 3.35106 / C. This compares very well with the experimentally determined value of 3.0106 / C [25]. The fiber damage was evaluated in terms of a decrease in Young’s modulus of the composite. In the 850 C sample after three cycles, a measured decrease in Young’s modulus from 155 to 85 GPa was observed. Thermal shock after one cycle resulted in no decrease in modulus, which once again, correlated well with microstructural observations. In thermal cycling of this composite, the microstructures were essentially crack-free [26]. The nondestructive nature of modulus measurement would seem to be more attractive than destructive testing, e.g., flexural testing after thermal shock [27]. It is important to note that flexural testing may not adequately quantify damage in thermal shocked materials, because of the combination of tensile, compressive, and shear stresses involved. Thus, while a predominantly tensile mode may operate in the undamaged composite, the presence of microcracks in the damaged state may reduce the interlaminar shear strength and the composite may fail in a shear mode [14]. Furthermore, shifting of the neutral axis in continuous fiber ceramic composites results in a significant enhancement of the strength [28,29]. It is interesting to note that microstructural damage and decrease in modulus are also very much affected by hold time prior to thermal shock. Blissett et al. [10] determined the thermal shock resistance of unidirectional and cross-plied Nicalon fiber reinforced CAS glass-ceramic composites. The composite specimens were allowed to age for one hour prior to shock. Their results indicated a strength decrease at intermediate temperatures followed by a perceived ‘‘recovery’’ in strength at Table 2 Elastic and thermal properties of hybrid matrix and Nicalon fiber Young’s modulus (GPa) Coefficient of thermal expansion (106 / C) Volume fraction Hybrid matrix 173–195 2.86 0.58 Nicalon fiber 192 4.0 0.42 Table 3 Debond and sliding stresses of Nicalon fibers in BMAS, as measured by fiber pushout testing Fiber diameter (mm) debond (MPa) sliding (MPa) As-received 17.92.9 16.4 5.1 12.44.6 Shocked, 500 C 23.21.9 12.72.6 8.71.9 Shocked, 950 C 21.42.4 13.26.8 9.13.9 N. Chawla et al. / Composites Science and Technology 61 (2001) 1923–1930 1927

1928 N. Chawla et al. Composites Science and Technology 61(2001)1923-1930 higher temperatures. The behavior was rationalized by in determining the macroscopic and microscopic the fact that at intermediate temperatures oxidation of the mechanical behavior of continuous fiber reinforced com- carbon interphase took place, while at higher tempera- posites. The interfacial properties of the composite were tures, oxidation was followed by the formation of a silica measured in the as-received condition as well as after a rich interface which contributed to higher strengths, single thermal shock by fiber pushout technique. Two although failure took place in a non-damage tolerant extreme thermal shock temperatures were chosen to com- manner. Multiple thermal shocks resulted in an increase in pare the interfacial properties to that of the as-received cracking and a reduction in modulus, which contributed composite. The stress versus displacement behavior dur to increased propensity for oxygen ingression into the ing a fiber pushout test of the as-received material is given fiber/matrix interface. In the present study, specimens in Fig. 6. Fig. 6 shows that during a pushout test, the were held at temperature for a much shorter period of fiber undergoes elastic loading and Poisson expansion time, resulting in less damage after thermal shock followed by progressive debonding along the fiber Reinforcement architecture also plays a role in thermal matrix interface. After complete debonding the fiber shock resistance. Wang et al. [30, 31] showed that in slides through the matrix, until the indentor contacts the Nicalon fiber reinforced SiC, with a carbon fiber coating, matrix. The debond shear stress, Id, and the frictional the unidirectional and cross-ply composites had lower shear stress, Tf, were determined by thermal shock resistance than woven materials This was attributed to the non-uniform distribution of the sic debond matrix in the unidirectional and cross-ply materials, pro- Ld cessed by chemical vapor infiltration. Quantifying the and Trt fiber/matrix interfacial properties is extremely important where r is the radius of the fiber t is the thickness of the specimen, Fdcbond is the peak load preceding the sharp load drop that is associated with complete debonding while Friction is the load obtained following complete debonding and is associated with pure fiber sliding(see Fig. 6). Values for complete interfacial debonding and sliding stresses for as-received and shocked samples from fiber pushout tests are given in Table 4. A sligh decrease in interfacial! shear strength(both debond and sliding stresses)was observed after thermal shock. The decrease in interfacial strengths was about the same for both 500 and 950C shocked samples, i. e, interfacial 0 6 strength for shock in this temperature regime seems to be independent of the temperature amplitude excursion Fig. 6. An example of a fiber pushout test of an as-received Nicalon SEM images of pushed-out fibers are shown in Fig. 7 reinforced BMAS composite. Fig. 7(a) shows a single fiber, demonstrating a smooth Fig. 7.(a) Some of the interfacial coating has been removed from the fiber during pushout testing. A smooth mating matrix surface remained during fiber pushout. (b)The lighter colored ring around several Nicalon fibers indicates the fiber/matrix debonding. The dark regions in the matrix represent the distributed glassy phase within the BMAs matrix

higher temperatures. The behavior was rationalized by the fact that at intermediate temperatures oxidation of the carbon interphase took place, while at higher temperatures, oxidation was followed by the formation of a silicarich interface which contributed to higher strengths, although failure took place in a non-damage tolerant manner. Multiple thermal shocks resulted in an increase in cracking and a reduction in modulus, which contributed to increased propensity for oxygen ingression into the fiber/matrix interface. In the present study, specimens were held at temperature for a much shorter period of time, resulting in less damage after thermal shock. Reinforcement architecture also plays a role in thermal shock resistance. Wang et al. [30,31] showed that in Nicalon fiber reinforced SiC, with a carbon fiber coating, the unidirectional and cross-ply composites had lower thermal shock resistance than woven materials. This was attributed to the non-uniform distribution of the SiC matrix in the unidirectional and cross-ply materials, processed by chemical vapor infiltration. Quantifying the fiber/matrix interfacial properties is extremely important in determining the macroscopic and microscopic mechanical behavior of continuous fiber reinforced composites. The interfacial properties of the composite were measured in the as-received condition as well as after a single thermal shock by fiber pushout technique. Two extreme thermal shock temperatures were chosen to compare the interfacial properties to that of the as-received composite. The stress versus displacement behavior during a fiber pushout test of the as-received material is given in Fig. 6. Fig. 6 shows that during a pushout test, the fiber undergoes elastic loading and Poisson expansion, followed by progressive debonding along the fiber/ matrix interface. After complete debonding the fiber slides through the matrix, until the indentor contacts the matrix. The debond shear stress, d, and the frictional shear stress, f, were determined by: d ¼ Fdebond 2rt and f ¼ Ffriction 2rt ð7Þ where r is the radius of the fiber, t is the thickness of the specimen, Fdebond is the peak load preceding the sharp load drop that is associated with complete debonding, while Ffriction is the load obtained following complete debonding and is associated with pure fiber sliding (see Fig. 6). Values for complete interfacial debonding and sliding stresses for as-received and shocked samples from fiber pushout tests are given in Table 4. A slight decrease in interfacia! shear strength (both debond and sliding stresses) was observed after thermal shock. The decrease in interfacial strengths was about the same for both 500 and 950 C shocked samples, i.e., interfacial strength for shock in this temperature regime seems to be independent of the temperature amplitude excursion. SEM images of pushed-out fibers are shown in Fig. 7. Fig. 7(a) shows a single fiber, demonstrating a smooth Fig. 6. An example of a fiber pushout test of an as-received Nicalon reinforced BMAS composite. Fig. 7. (a) Some of the interfacial coating has been removed from the fiber during pushout testing. A smooth mating matrix surface remained during fiber pushout. (b) The lighter colored ring around several Nicalon fibers indicates the fiber/matrix debonding. The dark regions in the matrix represent the distributed glassy phase within the BMAS matrix. 1928 N. Chawla et al. / Composites Science and Technology 61 (2001) 1923–1930

N Chawla et al. / Composites Science and Technology 61(2001)1923-1930 matrix surface after push out and showing no evidence Acknowledgements of an interfacial coating remaining on the matrix Fig. 7(b)shows a number of slightly raised fibers on a Funding for this work was provided by a Faculty polished cross section, indicating debonding of the fiber Research Grant from the University of Alabama at matrix interface Birmingham. The authors would also like to thank It is plausible that recession of the interfacial carbon Corning, Inc for providing the material layer, on the Nicalon fiber surface, takes place, as it is well known that carbon oxidizes above 450 oC. bocca. cini et al.[ 12] also measured the interfacial character- References istics in the as-received and thermal shocked condition of a Nicalon reinforced borosilicate glass composite. In [1 Chawla KK. Ceramic matrix composites. London: Chapman and their study, they also did not notice a significant change Hall.1993. in interfacial shear strength, although they also noticed 2] ature ceramic matrix composites. bington. Cambridge: Woodhead Publishing. 1993 an increase in scatter of the data after shock. The latter 3 Boccaccini AR, Pearce D, Janczak J, Beier W, Ponton CB was attributed to localized microcracking near the fiber/ tion of cyclic thermal shock behaviour of fibre rein- matrix interface region, which would clearly influence forced glass matrix composites using non-destructive force reso- the fiber /matrix interfacial shear strength at a particular nance technique. Mater Sci Technol 1997: 13: 852-8 site. Thus. it should be noted that interfacial character 4 Zawada L, Wetherhold R. The effects of thermal fatigue on a Sic fibre/aluminosilicate glass composite J Mater Sci 1991: 26: 648-54 ization by this technique is highly localized and may not 5 Boccaccini AR, Strutt A, Vecchio K, Mendoza D, Chawla KK capture the macroscopic nature of damage in this com Ponton CB et al. Behavior of Nicalon fiber-reinforced glass- posite, e.g., fiber cracking, since a very small section of matrix composites under thermal cycling conditions. Composites the material is being evaluated during fiber pushout [6 Chawla N. Effect of laminate stacking sequence on the high fre- quency fatigue behavior of SCS-6 fiber reinforced Si3 N4 matrix composites. Metall Mater Trans A 1997: 28A: 2423 4. Conclusions [7 Sun E, Lin HT, Brennan JJ. Intermediate temperature environ. mental effects on boron nitride- coated silicon carbide-fiber-rein. forced glass-ceramic composites. J Am Ceram Soc 1997: 80: 609-14 Thermal shock after three cycles induced a sig- [8 Gadkaree K. Whisker reinforcement of glass ceramics. J Mater nificant amount of damage in the form of micro- Sci1991:26:4845-54. tructural damage i.e., fiber cracking, which [9 Gadkaree K. Particulate-fibre-reinforced glass matrix hybrid composites. J Mater Sci 1992: 27: 3827-34 correlated well with the observed modulus decrease [10 Blissett MJ, Smith PA. Yeomans JA Flexural mechanical prop- Crack propagation through the matrix and crack erties of thermally treated unidirectional and cross-ply nicalor deflection at the fiber/matrix interface were also nforced calcium aluminosilicate composites. J Mater Sci 1998 observed. Thermal shock after one cycle only resul 33:4181-90. ted in negligible microstructural damage and the [1] Blissett MJ, Smith PA, Yeomans JA. Thermal shock behaviour of unidirectional silicon carbide fibre reinforced calcium alumi- modulus did not change significantly nosilicate. J Mater Sci 1997: 32- 317-25 The similarity in coefficient of thermal expansion [2]Boccaccini AR, Janczak-Rusch J, Pearce DH, Kern H.Assess between Sic whiskers and bMAS matrix resulted in ment of damage induced by thermal shock in SiC-fiber-reinforced a hybrid matrix with a lower CTE than that of the borosilicate glass composites. Comp Sci Technol 1999: 59: 105-12. fiber. This resulted in the fiber being loaded in ter [13 Shin D-w, Auh KH, Knowles KM. SiC fiber/borosilicate glass sion and the matrix in compression during thermal composite(Part 3). J Ceram Soc Japan 1995: 103: 533-40. [14 Kagawa Y, Kurosawa N, Kishi T. Thermal shock resistance of shock, resulting in fiber fracture as the dominant Sic fibre- reinforced borosilicate glass and lithium aluminosilicate damage mechanism, as opposed to matrix cracking matrix composites. J Mater Sci 1993: 28: 735-41 From a design standpoint, this is a desirable char- [15 Schneibel JH, Sabol SM, Morrison J, Ludeman E, Carmichae acteristic because the fiber has a higher strength and CA. Cyclic thermal shock resistance of several advanced ceramics mposites. J Am Ceram Soc 1998: 81: 1888- can sustain higher tensile loads than the matrix [16 Boccaccini AR, Ponton CB. Chawla KK. Development and Interfacial strength measurements, by fiber push healing of matrix microcracks in fibre reinforced glass matri out technique, showed a slight decrease in inter- t by internal friction. Mater Sci Eng 1998; A241:141-50. facial strength with thermal shock, inarputed ne [8] ASTM C 1259.94. Standard test method for dynamic Young's temperature excursion. This was attr to [7 Richerson Dw. Modern ceramic engineering. New York:Marcel recession of the carbon-rich layer present at surface of the Nicalon fiber. Interfacial character modulus shear modulus and Poissons ratio ization by fiber pushout may not be the most sui- cs by impulse excitation of vibration. 1994 ble technique for quantifying damage by thermal [19 Eldridge JI. Desktop fiber push-out apparatus. NASA TM shock because of the somewhat macroscopic nat- 105341. December 1991 20 Eldridge JI. The evolution of interfacial sliding stresses during ure of damage in this material cyclic push-in testing of C- and BN-coated Hi-Nicalon fiber

matrix surface after push out and showing no evidence of an interfacial coating remaining on the matrix. Fig. 7(b) shows a number of slightly raised fibers on a polished cross section, indicating debonding of the fiber/ matrix interface. It is plausible that recession of the interfacial carbon layer, on the Nicalon fiber surface, takes place, as it is well known that carbon oxidizes above 450 C. Boccaccini et al. [12] also measured the interfacial characteristics in the as-received and thermal shocked condition of a Nicalon reinforced borosilicate glass composite. In their study, they also did not notice a significant change in interfacial shear strength, although they also noticed an increase in scatter of the data after shock. The latter was attributed to localized microcracking near the fiber/ matrix interface region, which would clearly influence the fiber/matrix interfacial shear strength at a particular site. Thus, it should be noted that interfacial characterization by this technique is highly localized and may not capture the macroscopic nature of damage in this composite, e.g., fiber cracking, since a very small section of the material is being evaluated during fiber pushout. 4. Conclusions . Thermal shock after three cycles induced a significant amount of damage in the form of microstructural damage, i.e., fiber cracking, which correlated well with the observed modulus decrease. Crack propagation through the matrix and crack deflection at the fiber/matrix interface were also observed. Thermal shock after one cycle only resulted in negligible microstructural damage and the modulus did not change significantly. . The similarity in coefficient of thermal expansion between SiC whiskers and BMAS matrix, resulted in a hybrid matrix with a lower CTE than that of the fiber. This resulted in the fiber being loaded in tension and the matrix in compression during thermal shock, resulting in fiber fracture as the dominant damage mechanism, as opposed to matrix cracking. From a design standpoint, this is a desirable characteristic because the fiber has a higher strength and can sustain higher tensile loads than the matrix. . Interfacial strength measurements, by fiber pushout technique, showed a slight decrease in interfacial strength with thermal shock, independent of temperature excursion. This was attributed to recession of the carbon-rich layer present at the surface of the Nicalon fiber. Interfacial characterization by fiber pushout may not be the most suitable technique for quantifying damage by thermal shock because of the somewhat macroscopic nature of damage in this material. Acknowledgements Funding for this work was provided by a Faculty Research Grant from the University of Alabama at Birmingham. The authors would also like to thank Corning, Inc. for providing the material. References [1] Chawla KK. Ceramic matrix composites. London: Chapman and Hall, 1993. [2] Brennan JJ. High temperature ceramic matrix composites. Abington, Cambridge: Woodhead Publishing, 1993. [3] Boccaccini AR, Pearce D, Janczak J, Beier W, Ponton CB. Investigation of cyclic thermal shock behaviour of fibre reinforced glass matrix composites using non-destructive force resonance technique. Mater Sci Technol 1997;13:852–8. [4] Zawada L, Wetherhold R. The effects of thermal fatigue on a SiC fibre/aluminosilicate glass composite. J Mater Sci 1991;26:648–54. [5] Boccaccini AR, Strutt A, Vecchio K, Mendoza D, Chawla KK, Ponton CBet al. Behavior of Nicalon fiber-reinforced glassmatrix composites under thermal cycling conditions. Composites 1998;29A:1343–52. [6] Chawla N. Effect of laminate stacking sequence on the high frequency fatigue behavior of SCS-6 fiber reinforced Si3N4 matrix composites. Metall Mater Trans A 1997;28A:2423. [7] Sun E, Lin HT, Brennan JJ. Intermediate temperature environmental effects on boron nitride-coated silicon carbide-fiber-reinforced glass-ceramic composites. J Am Ceram Soc 1997;80:609–14. [8] Gadkaree K. Whisker reinforcement of glass ceramics. J Mater Sci 1991;26:4845–54. [9] Gadkaree K. Particulate-fibre-reinforced glass matrix hybrid composites. J Mater Sci 1992;27:3827–34. [10] Blissett MJ, Smith PA, Yeomans JA. Flexural mechanical properties of thermally treated unidirectional and cross-ply nicalonreinforced calcium aluminosilicate composites. J Mater Sci 1998; 33:4181–90. [11] Blissett MJ, Smith PA, Yeomans JA. Thermal shock behaviour of unidirectional silicon carbide fibre reinforced calcium aluminosilicate. J Mater Sci 1997;32:317–25. [12] Boccaccini AR, Janczak-Rusch J, Pearce DH, Kern H. Assessment of damage induced by thermal shock in SiC-fiber-reinforced borosilicate glass composites. Comp Sci Technol 1999;59:105–12. [13] Shin D-W, Auh KH, Knowles KM. SiC fiber/borosilicate glass composite (Part 3). J Ceram Soc Japan 1995;103:533–40. [14] Kagawa Y, Kurosawa N, Kishi T. Thermal shock resistance of SiC fibre-reinforced borosilicate glass and lithium aluminosilicate matrix composites. J Mater Sci 1993;28:735–41. [15] Schneibel JH, Sabol SM, Morrison J, Ludeman E, Carmichael CA. Cyclic thermal shock resistance of several advanced ceramics and ceramic composites. J Am Ceram Soc 1998;81:1888–92. [16] Boccaccini AR, Ponton CB, Chawla KK. Development and healing of matrix microcracks in fibre reinforced glass matrix composites: assessment by internal friction. Mater Sci Eng 1998; A241:141–50. [17] Richerson DW. Modern ceramic engineering. New York: Marcel Dekker, 1992 p. 767. [18] ASTM C 1259-94. Standard test method for dynamic Young’s modulus, shear modulus, and Poisson’s ratio for advanced ceramics by impulse excitation of vibration. 1994. [19] Eldridge JI. Desktop fiber push-out apparatus. NASA TM 105341, December 1991. [20] Eldridge JI. The evolution of interfacial sliding stresses during cyclic push-in testing of C- and BN-coated Hi-Nicalon fiberN. Chawla et al. / Composites Science and Technology 61 (2001) 1923–1930 1929

《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_Thermal-shock behavior of a Nicalon-fiber-reinforced hybrid glass-ceramic composite

《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_Thermal-shock behavior of a Nicalon-fiber-reinforced hybrid glass-ceramic composite