E噩≈S Journal of the European Ceramic Society 22(2002)2349-2356 www.elsevier.com/locate/jeurc Tensile fracture behavior of continuous sic fiber-reinforced SiC matrix composites at elevated temperatures and correlation to in Situ constituent properties Shuqi guo*,, Yutaka Kagawa Institute of Industrial Science, The University of Tokyo, 7-22-1, Roppongi, Minato-ku, Tokyo 106-8558, Japan Received 21 March 2001; received in revised form 20 December 2001: accepted 6 January 2002 Abstract The tensile fracture behavior and tensile mechanical properties of polymer infiltration pyrolysis(PIP)-processed two-dimensional plain-woven fabric carbon-coated NicalonM SiC fiber and BN-coated Hi-Nicalon"M SiC fiber-reinforced SiC matrix composites have been investigated. Tensile testing of the composites was carried out in air between 298 and 1400 K. In situ fiber strength and interface shear stress were determined by fracture mirror size and pulled-out fiber length measurements. For the Nicalon/C/Sic. tensile strength remained nearly constant up to 800 K, and while the strength dropped from 140 MPa at 800 K to 41 MPa at 1200 K, with weakest link failure mode. For the Hi-Nicalon/BN/SiC, the tensile strength increased slightly with increase in test tem- perature up to 1200 K; however, a large decrease in the strength was observed at 1400 K. In the case of the Hi-Nicalon/BN/ SiC, the fracture was governed by fiber bundle strength. The temperature dependence of tensile strength and fracture behavior of both omposites was attributed to change of the in situ constituent properties with temperature. C 2002 Elsevier Science Ltd. All rights Keywords: Composites; Fiber strength; Interface shear stress; Mechanical properties; Polymer infiltration pyrolysis; SiC/ SiC 1. Introduction The damage evolution of the composites usually includes two fundamental regimes: (i)at lower stresses Continuous fiber-reinforced ceramic matrix compo- matrix cracking originating from defects in the matrix sites(CFCCs) have become an important class of and followed by interface debonding between the fiber materials for structural applications at elevated tem- and matrix, and (ii) at higher stresses, occurrence of peratures because of their improved flaw tolerance, fiber damage and ultimate failure.6. 7 If matrix cracking large fracture resistance, and noncatastrophic mode of can reach the fully-saturated state prior to composite failure comparing with monolithic ceramic materials. failure, the subsequent deformation and failure are Among CFCCs, SiC fiber-reinforced SiC matrix com- dominated entirely by the fiber flaw population, show- posites have been studied extensively in recent years ing a noncatastrophic failure. 6, 7 However, if the fibers because the Sic fiber shows a potential for applications are sufficiently weak and interface bonding compara- at elevated temperatures. These studies demonstrated tively strong due to interface reaction in an oxidizing that the tensile mechanical behavior and properties of environment at high temperatures, composite failure is the composites strongly depend on the in situ fiber also possible in the regime of matrix cracking, showing strength characteristics, interface properties and matrix a catastrophic fracture similar to the fracture of a cracking stress, as well as the fabrication processes. I-5 monolithic ceramic 8,9 The transition from non- catastrophic to catastrophic fracture depends on in situ constituent properties. Thus, it is important to under 81-0298-51-3613. stand this transition to clarify tensile fracture behavior go. jp(S. Guo). of the composites and to allow correlation with the in Materials Science. I-I Namiki. Tsukuba situ constituent properties. In situ constituent properties of the composites such as fiber strength, interface shear 0955-2219/02/S. see front matter C 2002 Elsevier Science Ltd. All rights reserved. PII:S0955-2219(02)00028-6
Tensile fracture behavior of continuous SiC fiber-reinforced SiC matrix composites at elevated temperatures and correlation to in situ constituent properties Shuqi Guo*,1, Yutaka Kagawa Institute of Industrial Science, The University of Tokyo, 7-22-1, Roppongi, Minato-ku, Tokyo 106-8558, Japan Received 21 March 2001; received in revised form 20 December 2001; accepted 6 January 2002 Abstract The tensile fracture behavior and tensile mechanical properties of polymer infiltration pyrolysis (PIP)-processed two-dimensional plain-woven fabric carbon-coated NicalonTM SiC fiber and BN-coated Hi-NicalonTM SiC fiber-reinforced SiC matrix composites have been investigated. Tensile testing of the composites was carried out in air between 298 and 1400 K. In situ fiber strength and interface shear stress were determined by fracture mirror size and pulled-out fiber length measurements. For the Nicalon/C/SiC, tensile strength remained nearly constant up to 800 K, and while the strength dropped from 140 MPa at 800 K to 41 MPa at 1200 K, with weakest link failure mode. For the Hi-Nicalon/BN/SiC, the tensile strength increased slightly with increase in test temperature up to 1200 K; however, a large decrease in the strength was observed at 1400 K. In the case of the Hi-Nicalon/BN/SiC, the fracture was governed by fiber bundle strength. The temperature dependence of tensile strength and fracture behavior of both composites was attributed to change of the in situ constituent properties with temperature. # 2002 Elsevier Science Ltd. All rights reserved. Keywords: Composites; Fiber strength; Interface shear stress; Mechanical properties; Polymer infiltration pyrolysis; SiC/SiC 1. Introduction Continuous fiber-reinforced ceramic matrix composites (CFCCs) have become an important class of materials for structural applications at elevated temperatures because of their improved flaw tolerance, large fracture resistance, and noncatastrophic mode of failure comparing with monolithic ceramic materials. Among CFCCs, SiC fiber-reinforced SiC matrix composites have been studied extensively in recent years because the SiC fiber shows a potential for applications at elevated temperatures. These studies demonstrated that the tensile mechanical behavior and properties of the composites strongly depend on the in situ fiber strength characteristics, interface properties and matrix cracking stress, as well as the fabrication processes.15 The damage evolution of the composites usually includes two fundamental regimes: (i) at lower stresses, matrix cracking originating from defects in the matrix and followed by interface debonding between the fiber and matrix, and (ii) at higher stresses, occurrence of fiber damage and ultimate failure.6,7 If matrix cracking can reach the fully-saturated state prior to composite failure, the subsequent deformation and failure are dominated entirely by the fiber flaw population, showing a noncatastrophic failure.6,7 However, if the fibers are sufficiently weak and interface bonding comparatively strong due to interface reaction in an oxidizing environment at high temperatures, composite failure is also possible in the regime of matrix cracking, showing a catastrophic fracture similar to the fracture of a monolithic ceramic.8,9 The transition from noncatastrophic to catastrophic fracture depends on in situ constituent properties. Thus, it is important to understand this transition to clarify tensile fracture behavior of the composites and to allow correlation with the in situ constituent properties. In situ constituent properties of the composites such as fiber strength, interface shear 0955-2219/02/$ - see front matter # 2002 Elsevier Science Ltd. All rights reserved. PII: S0955-2219(02)00028-6 Journal of the European Ceramic Society 22 (2002) 2349–2356 www.elsevier.com/locate/jeurceramsoc * Corresponding author. Fax: +81-0298-51-3613. E-mail address: guo.shuqi@nims.go.jp (S. Guo). 1 National Institute for Materials Science, 1-1 Namiki, Tsukuba, Ibaraki 305-0044, Japan
S Guo, Y. Kagawa/Journal of the European Ceramic Society 22(2002)2349-2356 stress and matrix cracking stress, can be determined by The polycarbosilane had a melting point of 514 K,an a tensile test and fractographic analysis and measure- average molecular weight of 2410, and its chemical ments. 0,I Thus, it is possible to evaluate correlation of composition was: 60 mass% Si, 40 mass% C,<l e mechanical behavior to in situ constituent proper- mass%O. The five infiltrated pre-preg sheets were ties. Although the correlation of tensile fracture beha- stacked and pressed, and the stacked sheets were then vior to in situ constituent properties has been rep for precursor. The composite precursor was pyrolyzed at in CVI-processed composites, 3. 2 this correlation PIP-processed composites is not well known. In the 1273K in a high purity N2 atmosphere. Approximately present study, tensile testing of the two PIP-processed 40 vol. of pores existed in the pyrolyzed composite SiC fiber-reinforced SiC matrix composites was carried due to the low yield weight of the polycarbosilane. To out at room and elevated temperatures. The matrix further reduce the porosity of the composite, a multiple racking stress, in situ fiber strength and interface shear PIP process was applied. After 10 PIP cycles, the total stress were obtained. The correlation of the tensile frac- fiber volume fraction, f, and the total porosity of the ture behavior of the composites to in situ constituent composite were 0.28 and s0.09, respectively. Here- properties was discusse after. the NicalonTM Sic fiber and Hi-Nicalon M Sic fiber-reinforced SiC matrix composites fabricated are denoted as Nicalon/ C/SiC and Hi-Nicalon/BN/ SiC 2. Experimental procedure respectively 2. 1. Composite materials 2. 2. Tensile test The composite materials used in the present study The composite panels were cut into a dog-bone type were 2D plain-woven fabric SiC fiber-SiC matrix com- tensile test specimen with the long axis parallel to one of posites fabricated by the Pip process. To compare the the fiber axis directions. The shape and dimensions of effect of fiber strength and interface properties on the the specimen for monotonic tension testing are shown in tensile fracture behavior of the composites, NicalonM Fig. 1. Quasi-static monotonic tensile testing was car- and Hi-NicalonM SiC fibers(Nippon Carbon Co Ltd ried out using a servo-hydraulic MTs 808 testing system Tokyo, Japan)in which the surfaces were respectively (MTS System Co., MI, USA)with a constant crosshead coated with 0.04 um amorphous carbon and 0.4 um displacement rate of 0.5 mm/ min in ambient air at room bn by chemical vapor deposition(CVD) were used as temperature (298 K), 800, 1200 and 1400 K. Three reinforcements. The typical properties and chemical composite specimens were used for each measurement composition of the two fibers are listed in Table 1. The An electric furnace attached to the mts testing system coated SiC fibers were formed a plain-woven fabric provided the heating Axial strain was measured directly sheet, with 16x 16 numbers of fiber per inch. The fabric from the gauge length of the specimen by using a con- sheets averaging 150x 150 mm in size were cut from the tact extensometer (MTS Model 632.59, MTS System formed plane-woven sheet. The cut fabric sheets were Co., MI). Before the loading, the specimen was heated infiltrated with a polycarbosilane solution containing a fine B-SiC powder. The fine B-Sic powder had an aver- age diameter of N4 um and its addition effectively (A)Top View 90° Bundle reduced pore content in the SiC matrix after pyrolysis 4.5 0° Bundle Copper T able i Typical properties and chemical composition of the two fibers Fibres Nicalon Hi. Nicalon TM 25 25 Fibre properties Tensile strength (GPa) Youngs modulus (GPa) Elongation (B) Side View Average fibre radius Number of fibres Copper Tab Chemical compositions Si (wt%)56.6 (wt.%)31.7 (wt%) Fig. 1. Shape and dimensions of the specimen for monotonic tension testing
stress and matrix cracking stress, can be determined by a tensile test and fractographic analysis and measurements.10,11 Thus, it is possible to evaluate correlation of the mechanical behavior to in situ constituent properties. Although the correlation of tensile fracture behavior to in situ constituent properties has been reported in CVI-processed composites,3,12 this correlation for PIP-processed composites is not well known. In the present study, tensile testing of the two PIP-processed SiC fiber-reinforced SiC matrix composites was carried out at room and elevated temperatures. The matrix cracking stress, in situ fiber strength and interface shear stress were obtained. The correlation of the tensile fracture behavior of the composites to in situ constituent properties was discussed. 2. Experimental procedure 2.1. Composite materials The composite materials used in the present study were 2D plain-woven fabric SiC fiber–SiC matrix composites fabricated by the PIP process. To compare the effect of fiber strength and interface properties on the tensile fracture behavior of the composites, NicalonTM and Hi-NicalonTM SiC fibers (Nippon Carbon Co. Ltd., Tokyo, Japan) in which the surfaces were respectively coated with 0.04 mm amorphous carbon and 0.4 mm BN by chemical vapor deposition (CVD) were used as reinforcements. The typical properties and chemical composition of the two fibers are listed in Table 1.13 The coated SiC fibers were formed a plain-woven fabric sheet, with 1616 numbers of fiber per inch. The fabric sheets averaging 150150 mm in size were cut from the formed plane-woven sheet. The cut fabric sheets were infiltrated with a polycarbosilane solution containing a fine b-SiC powder. The fine b-SiC powder had an average diameter of 4 mm and its addition effectively reduced pore content in the SiC matrix after pyrolysis.4,5 The polycarbosilane had a melting point of 514 K, an average molecular weight of 2410, and its chemical composition was: 60 mass% Si, 40 mass% C, <1 mass% O. The five infiltrated pre-preg sheets were stacked and pressed, and the stacked sheets were then cured at 523 Kin ambient air to obtain the composite precursor. The composite precursor was pyrolyzed at 1273 Kin a high purity N2 atmosphere. Approximately 40 vol.% of pores existed in the pyrolyzed composite due to the low yield weight of the polycarbosilane. To further reduce the porosity of the composite, a multiple PIP process was applied. After 10 PIP cycles, the total fiber volume fraction, f, and the total porosity of the composite were 0.28 and 0.09, respectively. Hereafter, the NicalonTM SiC fiber and Hi-NicalonTM SiC fiber-reinforced SiC matrix composites fabricated are denoted as Nicalon/C/SiC and Hi-Nicalon/BN/SiC, respectively. 2.2. Tensile test The composite panels were cut into a dog-bone type tensile test specimen with the long axis parallel to one of the fiber axis directions. The shape and dimensions of the specimen for monotonic tension testing are shown in Fig. 1. Quasi-static monotonic tensile testing was carried out using a servo-hydraulic MTS 808 testing system (MTS System Co., MI, USA) with a constant crosshead displacement rate of 0.5 mm/min in ambient air at room temperature (298 K), 800, 1200 and 1400 K. Three composite specimens were used for each measurement. An electric furnace attached to the MTS testing system provided the heating. Axial strain was measured directly from the gauge length of the specimen by using a contact extensometer (MTS Model 632.59, MTS System Co., MI). Before the loading, the specimen was heated Table 1 Typical properties and chemical composition of the two fibers Fibres NicalonTM SiC Hi-NicalonTM SiC Fibre properties Tensile strength (GPa) 3.0 2.8 Young’s modulus (GPa) 220 270 Elongation (%) 1.4 1.0 Average fibre radius (mm) 7 7 Number of fibres per bundles 500 500 Chemical compositions Si (wt.%) 56.6 62.4 C (wt.%) 31.7 37.1 O (wt%) 11.7 0.5 Fig. 1. Shape and dimensions of the specimen for monotonic tension testing. 2350 S. Guo, Y. Kagawa / Journal of the European Ceramic Society 22 (2002) 2349–2356
S Guo, Y. Kagawa/Journal of the European Ceramic Society 22(2002)2349-2356 to the test temperature at a rate of 50C/min and then response. The large nonlinear regime is observed and held for a10 min to give a uniform temperature dis- remains nearly constant up to 800 K, however, this tribution in the specimen. After the tension testing, the regime is sharply reduced at 1200 K and nearly dis fracture surfaces of the specimens were characterized appears, showing a brittle fracture similar to that of using an optical microscope and scanning electron monolithic ceramics. Fig. 3 shows the typical tensile microscope(SEM) stress-strain curves of the Hi-Nicalon/ BN/SiC at room and elevated temperatures. The curves show a tensile 2.3. In situ fiber strength and effective interface shear stress-strain behavior similar to that of the Nicalon/C/ stress SiC: the linear deformation regime and nonlinear deformation regime are present, but the nonlinear (i) In situ fiber strength: in situ fiber strength in the regime is larger for all the test temperatures compared composite, ofu, was estimated from the fracture mirror to the Nicalon/C/SiC, showing a larger fracture resis radius, Im, of a pulled-out fiber. 0 The tensile strength, tance, especially at and above 1200 K Ofu, of the fiber is given by: The Youngs modulus of the composite, Ec, was obtained from the initial slope of these curves. The apparent matrix cracking stress, omc, was determined at m the transition point from a linear response to a non linear response, because the deviation from linear beha- where Kr is the(mode I)fracture toughness of the fiber. vior is often attributed to initiation of the matrix The fracture toughness of the SiC fiber is reported to be cracking. 18, 19 However, studies have shown that matrix A1.0 MPa m/ 14 This value is assumed to be indepen dent of the test tempe rature dependence up to 1100 K has been observed for N 800K lonM SiC fiber. 14 Assuming that the fiber failure follows the weakest-link principle, the statistical distribution of 298K fiber strength is described according to Weibull's two- 100 parameter statistical distribution theory. 5 Failure probability of the fiber, which is associated with fiber strength, is obtained using the mean rank method (ii)Interface shear stress: interface shear stress, Ti, of Sic fiber-reinforced ceramic matrix composites is expressed using an average fiber pullout length, L 1200K 0.5 Fig. 2. Typical urves of the Nicalon/ C/SiC at room and where the Rr is the radius of fiber(7 um), i(m) is a nondimensional function determined from the statistics 300 of fiber failure properties and takes a value close to Hi-Nicalon/BN/SiC 1200K unity for m>3. The interface shear stress determined represents the load transverse capacity from the fiber to he matrix, reflecting the interface sliding resistance in G200 the debonding interface. 1400K 3. Results 3.. Tensile mechanical behavior (i Tensile stress-strain curves: Fig. 2 shows the typi- cal tensile stress-strain curves of the Nicalon/ C/SiC at room aI nd elevated temperatures. All the curves exhibit Strain, E(%) a linear response near the origin and a following gra- Fig 3. Typical monotonic tensile stress-strain curves of the Hi-Nica- dual decrease of the slope up to fracture, i.e. nonlinear lon/BN/SiC at room and elevated temperatures
to the test temperature at a rate of 50 C/min and then held for 10 min to give a uniform temperature distribution in the specimen. After the tension testing, the fracture surfaces of the specimens were characterized using an optical microscope and scanning electron microscope (SEM). 2.3. In situ fiber strength and effective interface shear stress (i) In situ fiber strength: in situ fiber strength in the composite, fu, was estimated from the fracture mirror radius, rm, of a pulled-out fiber.10 The tensile strength, fu, of the fiber is given by:10 fu ¼ 3:5Kf ffiffiffiffiffi rm p ð1Þ where Kf is the (mode I) fracture toughness of the fiber. The fracture toughness of the SiC fiber is reported to be 1.0 MPa m1/2. 14 This value is assumed to be independent of the test temperature because only a slight dependence up to 1100 Khas been observed for NicalonTM SiC fiber.14 Assuming that the fiber failure follows the weakest-link principle, the statistical distribution of fiber strength is described according to Weibull’s twoparameter statistical distribution theory.15 Failure probability of the fiber, which is associated with fiber strength, is obtained using the mean rank method. (ii) Interface shear stress: interface shear stress, i, of SiC fiber-reinforced ceramic matrix composites is expressed using an average fiber pullout length, L, as14,16 i ¼ lðmÞRffu 4L ð2Þ where the Rf is the radius of fiber ( 7 mm), l(m) is a nondimensional function determined from the statistics of fiber failure properties and takes a value close to unity for m>3.17 The interface shear stress determined represents the load transverse capacity from the fiber to the matrix, reflecting the interface sliding resistance in the debonding interface. 3. Results 3.1. Tensile mechanical behavior (i) Tensile stress–strain curves: Fig. 2 shows the typical tensile stress–strain curves of the Nicalon/C/SiC at room and elevated temperatures. All the curves exhibit a linear response near the origin and a following gradual decrease of the slope up to fracture, i.e. nonlinear response. The large nonlinear regime is observed and remains nearly constant up to 800 K; however, this regime is sharply reduced at 1200 Kand nearly disappears, showing a brittle fracture similar to that of monolithic ceramics. Fig. 3 shows the typical tensile stress–strain curves of the Hi-Nicalon/BN/SiC at room and elevated temperatures. The curves show a tensile stress–strain behavior similar to that of the Nicalon/C/ SiC; the linear deformation regime and nonlinear deformation regime are present, but the nonlinear regime is larger for all the test temperatures compared to the Nicalon/C/SiC, showing a larger fracture resistance, especially at and above 1200 K. The Young’s modulus of the composite, Ec, was obtained from the initial slope of these curves. The apparent matrix cracking stress, mc, was determined at the transition point from a linear response to a nonlinear response, because the deviation from linear behavior is often attributed to initiation of the matrix cracking.18,19 However, studies have shown that matrix Fig. 2. Typical monotonic tensile stress–strain curves of the Nicalon/ C/SiC at room and elevated temperatures. Fig. 3. Typical monotonic tensile stress–strain curves of the Hi-Nicalon/BN/SiC at room and elevated temperatures. S. Guo, Y. Kagawa / Journal of the European Ceramic Society 22 (2002) 2349–2356 2351
S Guo, Y. Kagawa/Journal of the European Ceramic Society 22(2002)2349-2356 cracking usually initiates well below this stress 20.21 tested at room and elevated temperatures. The pulled- Thus, in the present study, the apparent matrix cracking out fibers are observed in the composite specimens stress obtained should be higher than the true matrix fractured at and below 800 K, however, this fiber pull- cracking stress. The matrix cracking behavior clearly out behavior is not seen in the composite tested at 1200 was observed in the polished longitudinal cross-section K, showing a brittle fracture fashion. Fig. 6 shows the of the composite specimens tested at room and elevated macroscopic fracture appearances of the Hi-Nicalon temperatures(Fig 4). The transverse matrix cracks are BN SiC tested at room and elevated temperatures. Dif- formed through the transverse fiber bundle, and some of fering from the Nicalon /C/SiC, the room temperature them are arrested at the interface between the long- fracture behavior of the Hi-Nicalon/BN/SiC shows a itudinal and transverse fiber bundles fracture path which is jagged and stepped across the The average values of Ec and ome that were calculated thickness, and there is extensive fiber and fiber tow pull from the duplicate tests at each temperature are shown out. Although the fracture surface became smoother in Table 2, together with the tensile strength, OTs, and with increasing test temperature, the pulled-out fibers le strain to failure, Ec. For the Nicalon/C/SiC, the ten- are observed over all test temperatures sile mechanical properties remain nearly constant from Fig. 7 shows SEM micrographs of the fracture sur- 298 to 800 K and that they degrade sharply by 1200 K, faces of both the composites. Although all fracture sur in particular, the tensile strength and strain to failure faces of the composites tested at room and hig dropped from 140 MPa and 0. 38% at 800 K to 41 MPa temperatures showed a fibrous fracture surface, only and 0.075% at 1200 K. For the Hi-Nicalon/ BN/SiC, on few and short pulled-out fibers are observed in the the other hand, the mechanical properties such as Ec, fracture surface of the Nicalon/C/SiC tested at 1200 K ome, Ec and ars remain nearly constant up to 1200 K This is probably attributed to the silica(Sio2) formation and begin to degrade at 1400 K; they are much higher at the interface by oxidation of both the matrix and than those of the Nicalon/C/SiC at all test temperatures fiber, because of air penetration into the gaps at the interface resulting form the elimination of the C-coating (ii) Fracture surface observations: Fig. 5 shows the layer by oxidation above 700 K.,3 The pulled-out macroscopic fracture appearance of the Nicalon/ C/Sic length of fiber was measured using the method reported elsewhere. 3 For the Nicalon/C/SiC, the average pulled ut fiber length is≈50mat298K,≈70pmat800K and a20 um at 1200 K. For the Hi-Nicalon/BN/SIC. the pulled-out fiber length is a300 um at 298 K, N250 mat1200Kand≈190μmatl400K. It is clear that he pulled-out fiber length of the Hi-Nicalon/ BN/SiC is much larger than that of the Nicalon/ C/Sic at room and elevated temperatures. Moreover, for the Nicalon/ C/Sic tested at and below 800 K there are regions in all the bundles in which the fiber -fracture locations are essentially coplanar with one another compared to the Hi-Nicalon/BN/SIC. 10m 3.2. In situ fiber strength Fig 4. An example of optical photographs of polished longitudinal The in situ fiber strength characteristics of the Nica- cross-section of the two composites after monotonic tension testing. lon TM Sic fiber and the hi-Nicalon TM Sic fiber were showing matrix cracking(T=298 K, Hi-Nicalon/BN/SiC obtained using Eq (I)and the results are summarized in Table 2 Tensile experimental results of both the Nicalon/C/SiC and Hi-Nicalon/ BN/SIC Composite temperature T(K) modulus E(GPa) cracking stress me(MPa)strength oTs(MPa) failure e(%)strength oTs Measured tensile Strain to Predicted tensi Nicalon/C/SiC 58±5 65±8 136±19 0.42士 55±4 5±5 0.38士 17: 1200 33±3 4l±5 0.075±0.01512 Hi-Nicalon/BN/SIC 298 80±5 75±9 226±l1 0.84±0.12223 76士4 0±7 237士6 0.9士0.06228 0±5 197±15 0.68±0.09209
cracking usually initiates well below this stress.20,21 Thus, in the present study, the apparent matrix cracking stress obtained should be higher than the true matrix cracking stress. The matrix cracking behavior clearly was observed in the polished longitudinal cross-section of the composite specimens tested at room and elevated temperatures (Fig. 4). The transverse matrix cracks are formed through the transverse fiber bundle, and some of them are arrested at the interface between the longitudinal and transverse fiber bundles. The average values of Ec and mc that were calculated from the duplicate tests at each temperature are shown in Table 2, together with the tensile strength, TS, and the strain to failure, c. For the Nicalon/C/SiC, the tensile mechanical properties remain nearly constant from 298 to 800 Kand that they degrade sharply by 1200 K, in particular, the tensile strength and strain to failure dropped from 140 MPa and 0.38% at 800 Kto 41 MPa and 0.075% at 1200 K. For the Hi-Nicalon/BN/SiC, on the other hand, the mechanical properties such as Ec, mc, c and TS remain nearly constant up to 1200 K and begin to degrade at 1400 K; they are much higher than those of the Nicalon/C/SiC at all test temperatures. (ii) Fracture surface observations: Fig. 5 shows the macroscopic fracture appearance of the Nicalon/C/SiC tested at room and elevated temperatures. The pulledout fibers are observed in the composite specimens fractured at and below 800 K, however, this fiber pullout behavior is not seen in the composite tested at 1200 K, showing a brittle fracture fashion. Fig. 6 shows the macroscopic fracture appearances of the Hi-Nicalon/ BN/SiC tested at room and elevated temperatures. Differing from the Nicalon/C/SiC, the room temperature fracture behavior of the Hi-Nicalon/BN/SiC shows a fracture path which is jagged and stepped across the thickness, and there is extensive fiber and fiber tow pullout. Although the fracture surface became smoother with increasing test temperature, the pulled-out fibers are observed over all test temperatures. Fig. 7 shows SEM micrographs of the fracture surfaces of both the composites. Although all fracture surfaces of the composites tested at room and high temperatures showed a fibrous fracture surface, only few and short pulled-out fibers are observed in the fracture surface of the Nicalon/C/SiC tested at 1200 K. This is probably attributed to the silica (SiO2) formation at the interface by oxidation of both the matrix and fiber, because of air penetration into the gaps at the interface resulting form the elimination of the C-coating layer by oxidation above 700 K.22,23 The pulled-out length of fiber was measured using the method reported elsewhere.3 For the Nicalon/C/SiC, the average pulledout fiber length is 50 mm at 298 K, 70 mm at 800 K and 20 mm at 1200 K. For the Hi-Nicalon/BN/SiC, the pulled-out fiber length is 300 mm at 298 K , 250 mm at 1200 Kand 190 mm at 1400 K. It is clear that the pulled-out fiber length of the Hi-Nicalon/BN/SiC is much larger than that of the Nicalon/C/SiC at room and elevated temperatures. Moreover, for the Nicalon/ C/SiC tested at and below 800 Kthere are regions in all the bundles in which the fiber-fracture locations are essentially coplanar with one another compared to the Hi-Nicalon/BN/SiC. 3.2. In situ fiber strength The in situ fiber strength characteristics of the NicalonTM SiC fiber and the Hi-NicalonTM SiC fiber were obtained using Eq. (1) and the results are summarized in Fig. 4. An example of optical photographs of polished longitudinal cross-section of the two composites after monotonic tension testing, showing matrix cracking (T=298 K, Hi-Nicalon/BN/SiC). Table 2 Tensile experimental results of both the Nicalon/C/SiC and Hi-Nicalon/BN/SiC Composite materials Test temperature T (K) Young’s modulus Ec (GPa) Apparent matrix cracking stress mc (MPa) Measured tensile strength TS (MPa) Strain to failure c (%) Predicted tensile strength TS Nicalon/C/SiC 298 58 5 65 8 136 19 0.42 0.1 177 800 55 4 55 5 140 12 0.38 0.08 175 1200 49 4 33 3 41 5 0.075 0.015 127 Hi-Nicalon/BN/SiC 298 80 5 75 9 226 11 0.84 0.12 223 1200 76 4 70 7 237 6 0.9 0.06 228 1400 60 3 50 5 197 15 0.68 0.09 209 2352 S. Guo, Y. Kagawa / Journal of the European Ceramic Society 22 (2002) 2349–2356
S Guo, Y. Kagawa/Journal of the European Ceramic Society 22(2002)2349-2356 2353 罐翻 0.1mm 0.1mm (b) 0.1mm 0.1mm Fig. 5. Macroscopic fracture appearance of the Nicalon/C/SiC tested Fig. 6. Macroscopic fracture appearance of the Hi-Nicalon/BN/SiC at(a)298K,(b)800Kand(c)1200K tested at(a) 298 K,(b)1200 K and(c)1400 K. Table 3. The in situ strength of Nicalon TM SiC fiber is elevated temperature are expected for the carbon 1765 MPa at 298 K and 1705 MPa at 800 K; however, it coated NicalonTM fibers investigated in the present decreases to 1235 MPa at 1200 K. The degradation of study, especially microstructural and stoichimetric fiber strength due to exposure to high temperatures is changes of the fiber. On the other hand, the in situ already documented in the literature. 24,25 The fiber strength of Hi-Nicalon M SiC fiber remains nearly the strength decreased m 30 and 70%, respectively, after same value at the temperatures of 298 and 1200 K, and exposure at 1273 and 1573 K for 12 h in wet-air atmo- it begins to decrease above 1200 K. The fiber strength at sphere, 4 and the tensile strength drops from 2000 MPa 1400 K is lower by N11%, compared to that at room at room temperature to 1000 MPa at 1573 K in air. 5 temperature. Strength decrease of the Hi-NicalonTM This reduction is explained by the microstructural and SiC fiber at high temperature has been documented in stoichiometric changes and void formation in the fiber the literature. 26 Degradation in the strength of the fiber at higher temperatures. Similar changes, although not above the temperature of 1573 K in ambient air as severe because of a short-term heat exposure at observed and the major reason given for this was
Table 3. The in situ strength of NicalonTM SiC fiber is 1765 MPa at 298 Kand 1705 MPa at 800 K; however, it decreases to 1235 MPa at 1200 K. The degradation of fiber strength due to exposure to high temperatures is already documented in the literature.24,25 The fiber strength decreased 30 and 70%, respectively, after exposure at 1273 and 1573 Kfor 12 h in wet-air atmosphere,24 and the tensile strength drops from 2000 MPa at room temperature to 1000 MPa at 1573 Kin air.25 This reduction is explained by the microstructural and stoichiometric changes and void formation in the fiber at higher temperatures. Similar changes, although not as severe because of a short-term heat exposure at elevated temperature, are expected for the carboncoated NicalonTM fibers investigated in the present study, especially microstructural and stoichimetric changes of the fiber. On the other hand, the in situ strength of Hi-NicalonTM SiC fiber remains nearly the same value at the temperatures of 298 and 1200 K, and it begins to decrease above 1200 K. The fiber strength at 1400 Kis lower by 11%, compared to that at room temperature. Strength decrease of the Hi-NicalonTM SiC fiber at high temperature has been documented in the literature.26 Degradation in the strength of the fiber above the temperature of 1573 Kin ambient air was observed and the major reason given for this was grain Fig. 5. Macroscopic fracture appearance of the Nicalon/C/SiC tested at (a) 298 K, (b) 800 K and (c) 1200 K. Fig. 6. Macroscopic fracture appearance of the Hi-Nicalon/BN/SiC tested at (a) 298 K, (b) 1200 K and (c) 1400 K. S. Guo, Y. Kagawa / Journal of the European Ceramic Society 22 (2002) 2349–2356 2353