COMPOSITES SCIENCE AND TECHNOLOGY ELSEVIER Composites Science and Technology 60(2000)1067-1076 The embrittlement of Nicalon/alumina composites at intermediate and elevated temperatures J.A. Celemin . J. LLorca Department of Materials Science, Polytechnic University of Madrid, E. T.S. de Ingenieros de Caminos, 28040 Madrid, Spain Received 13 April 1999: received in revised form 2 November 1999; accepted 22 December 1999 The strength and toughness of a 2-D woven Nicalon/Al,O3-matrix composite were measured at ambient, intermediate(800C) and elevated (1000-1200C)temperatures. The co tes exhibited non-linear behavior over the whole temperature range but their mechanical properties were significantly degraded at 800C and above. The in situ fiber strength and the interfacial sliding resistance were determined through quantitative microscopy techniques and they were used to predict the composite properties by means of using the appropriate micromechanical models. Comparison of the model and the experiments, together with the fracto- graphic observations, led to the conclusion that the strength reduction was caused by localized interface oxidation at 800 C and by the degradation of the fiber strength at 1000C and above. The decrease in the composite fracture energy was mainly induced by a transition in the fracture mode, which changed from the development of a diffuse damage zone with multiple matrix cracks at 25C to the propagation of a single dominant crack at 800C and above. C 2000 Elsevier Science Ltd. All rights reserved Keywords: Ceramic-matrix composites: Embrittlement; Mechanical properties; Fracture; Strength 1. Introduction (see, for instance, [7-9), aimed at elucidating the domi nant mechanisms of high-temperature degradation in It is now well established that tough, damage-tolerant these novel composites fiber-reinforced ceramics(FRCs) can be obtained by These investigations have pointed to interface oxida- appropriate design of the fiber/matrix interface. If the tion and fiber degradation as the main causes of the matrix is weakly bonded to the fibers, the matrix cracks reduction in strength and toughness at elevated tem- bifurcate at the interface and propagate along it. The perature [10]. In addition, several FRCs were identified cracks propagate through the matrix upon further whose mechanical properties exhibited a minimum loading, leaving the intact fibers in the crack wake. when tested (or exposed prior to testing)at intermediate These processes activate several mechanisms of energy (500-1000C)rather than at elevated temperatures [11 dissipation, such as multiple matrix cracking, crack 13]. This is important from the application standpoint deflection, and fiber failure and pull-out, which lead to because the structural integrity of the components non-linear behavior and to a remarkable improvement should be guaranteed over the whole temperature range in the fracture toughness [1-3. These FRCs were tar- of operation. However, the differences and similarities in geted for structural applications at elevated tempera- the embrittlement mechanisms between intermediate tures(>1000C), where the nickel-based superalloys (600-900 C)and elevated temperature(>1000 C)are become inadequate as they approach their melting not yet well established, and this is partially a result of point. It was soon found, however, that FRC often the lack of systematic studies on the mechanical prop experienced severe embrittlement when they were tested erties of FRCs as a function of temperature. In order to at elevated temperatures(>1000oC)in oxidizing envir- contribute to this research effort, the mechanical prop onments [4-6]. This prompted a huge amount of work erties of an alumina matrix reinforced with Nicalon Sic on the high-temperature mechanical behavior of FRCs fibers were measured at 25, 800, 1000 and 1200 C. The critical microstructural parameters which control the composite properties (interfacial sliding resistance and responding author fiber strength) were estimated as a function of tempera Current address: Universidad Pontificia Comillas de madrid ture through quantitative microscopy techniques. They 0266-3538/00/S. see front matter C 2000 Elsevier Science Ltd. All rights reserved PII:S0266-3538(00)00007-5
The embrittlement of Nicalon/alumina composites at intermediate and elevated temperatures J.A. CelemõÂn 1 , J. LLorca * Department of Materials Science, Polytechnic University of Madrid, E. T. S. de Ingenieros de Caminos, 28040 Madrid, Spain Received 13 April 1999; received in revised form 2 November 1999; accepted 22 December 1999 Abstract The strength and toughness of a 2-D woven Nicalon/Al2O3-matrix composite were measured at ambient, intermediate (800C), and elevated (1000±1200C) temperatures. The composites exhibited non-linear behavior over the whole temperature range but their mechanical properties were signi®cantly degraded at 800C and above. The in situ ®ber strength and the interfacial sliding resistance were determined through quantitative microscopy techniques and they were used to predict the composite properties by means of using the appropriate micromechanical models. Comparison of the model and the experiments, together with the fractographic observations, led to the conclusion that the strength reduction was caused by localized interface oxidation at 800C and by the degradation of the ®ber strength at 1000C and above. The decrease in the composite fracture energy was mainly induced by a transition in the fracture mode, which changed from the development of a diuse damage zone with multiple matrix cracks at 25C to the propagation of a single dominant crack at 800C and above. # 2000 Elsevier Science Ltd. All rights reserved. Keywords: Ceramic-matrix composites; Embrittlement; Mechanical properties; Fracture; Strength 1. Introduction It is now well established that tough, damage-tolerant ®ber-reinforced ceramics (FRCs) can be obtained by appropriate design of the ®ber/matrix interface. If the matrix is weakly bonded to the ®bers, the matrix cracks bifurcate at the interface and propagate along it. The cracks propagate through the matrix upon further loading, leaving the intact ®bers in the crack wake. These processes activate several mechanisms of energy dissipation, such as multiple matrix cracking, crack de¯ection, and ®ber failure and pull-out, which lead to non-linear behavior and to a remarkable improvement in the fracture toughness [1±3]. These FRCs were targeted for structural applications at elevated temperatures (>1000C), where the nickel-based superalloys become inadequate as they approach their melting point. It was soon found, however, that FRC often experienced severe embrittlement when they were tested at elevated temperatures (>1000C) in oxidizing environments [4±6]. This prompted a huge amount of work on the high-temperature mechanical behavior of FRCs (see, for instance, [7±9]), aimed at elucidating the dominant mechanisms of high-temperature degradation in these novel composites. These investigations have pointed to interface oxidation and ®ber degradation as the main causes of the reduction in strength and toughness at elevated temperature [10]. In addition, several FRCs were identi®ed, whose mechanical properties exhibited a minimum when tested (or exposed prior to testing) at intermediate (500±1000C) rather than at elevated temperatures [11± 13]. This is important from the application standpoint because the structural integrity of the components should be guaranteed over the whole temperature range of operation. However, the dierences and similarities in the embrittlement mechanisms between intermediate (600±900C) and elevated temperature (>1000C) are not yet well established, and this is partially a result of the lack of systematic studies on the mechanical properties of FRCs as a function of temperature. In order to contribute to this research eort, the mechanical properties of an alumina matrix reinforced with Nicalon SiC ®bers were measured at 25, 800, 1000 and 1200C. The critical microstructural parameters which control the composite properties (interfacial sliding resistance and ®ber strength) were estimated as a function of temperature through quantitative microscopy techniques. They 0266-3538/00/$ - see front matter # 2000 Elsevier Science Ltd. All rights reserved. PII: S0266-3538(00)00007-5 Composites Science and Technology 60 (2000) 1067±1076 * Corresponding author. 1 Current address: Universidad Ponti®cia Comillas de Madrid
J.A. Celemin, J. LLorca/ Composites Science and Technology 60(2000)1067-1076 were used to predict the composite strength and toug men for an accurate control of the temperature. The ness using the available micromechanical data for FRCs heating rate was 12 C per min and each specimen was with a weak interface. The comparison between model held at the test temperature for I h prior to testing predictions and experimental results, together with the ensure uniform distribution of the temperature. Both microstructural analysis of the deformation and failure tensile and fracture tests were performed under stroke processes, shed more light on embrittlement mechan- control at a crosshead speed of 50 um /min The load (P) isms which control the mechancial behavior of this and the crosshead displacement of the testing machine material as a function of temperature with respect to the frame(v) were monitored during the tests, the latter through a lvdt transducer outside the furnace. In addition. the distance between two small 2. Materials and experimental techniques pins glued in the central region of the tensile specimens was measured through a laser extensometer. This The composite material was supplied by Dupont- extensometer was also used during the fracture tests to Lanxide Corporation(Newark, DE)in the form of rec- measure the crack mouth opening displacement angular plates of 3 mm nominal thickness. The fiber (CMOD)obtained from the distance between two alu- preform was manufactured by stacking several layers of mina pins glued symmetrically to the notch mouth. bi-directional (0 /90)Nicalon &-harness satin-weave Once broken, the fracture surfaces were sputtered fabric(Nippon Carbon, Tokyo, Japan). The average with Au-Pd for 3 min before being observed in the fiber radius, as determined by quantitative microscopy, scanning electron microscope. In addition, the speci was 7. 2 um. The fibers were coated by chemical var mens were sliced far away from the fracture surface with deposition with a thin layer of BN (approximately 100 a low speed diamond saw, and the longitudinal sections nm)and then with a thicker layer of Sic(3 um) onto were polished successively on diamond cloths of 40, 9, 3 the BN. The alumina matrix was infiltrated by using and 1 um grain size and finally on alumina of 0.3 um a direct metal oxidation process. The preform was grain size. They were cleaned for 30 min by ultrasound brought into contact with molten aluminum in air at in acetone to remove the alumina from polishing, and studied in the optical form a matrix of porous Al,O3 which grew into the preform. The residual aluminum was finally removed from the composite material using a proprietary techni- 3. Experimental results que. The volume fraction of fibers in the composite was 37%, while the Sic fiber coating occupied another 37%. 3.1. Mechanical properties The Al2O3 matrix comprised 18% of the composite volume and the rest(8%)was porosity. More details of The initial composite response during the tensile tests the manufacturing process and the microstructure can was linear. It was followed by a pronounced knee in the be found elsewhere [14-16 Tensile and notched-beam specimens were machined rom the plates. The tensile specimens had a dog-bone shape and were designed according to the specifications given by French and Wiederhorn [17]. The central part Thermocouples of the specimen had uniform width of 4 mm and a length of 15 mm. Load was applied through two cera- mic pins which were introduced into circular holes Hinge machined in the specimen heads. The pins were attached to two alumina rods connected, respectively to the actuator and to the load cell of the mechanical testing machine. Two bi-directional hinges were inserted in the Heating elements oad train to avoid bending stresses during the tests. The Lase experimental set-up is shown in Fig. 1. The fracture tests were carried out by three-point bending of notched prismatic bars. The loading span was 50 mm and the bar depth 10 mm. A notch of around 2 mm in length and 150 um in radius was introduced with a thin dia- Alumina rod mond wire LVDT The specimen and the loading fixture were placed in a urnace for the elevated temperature tests. (see Fig. 1) Two B-type thermocouples were attached to the speci- Fig. 1. Experimental set-up for the tensile tests at elevated temperature
were used to predict the composite strength and toughness using the available micromechanical data for FRCs with a weak interface. The comparison between model predictions and experimental results, together with the microstructural analysis of the deformation and failure processes, shed more light on embrittlement mechanisms which control the mechancial behavior of this material as a function of temperature. 2. Materials and experimental techniques The composite material was supplied by DupontLanxide Corporation (Newark, DE) in the form of rectangular plates of 3 mm nominal thickness. The ®ber preform was manufactured by stacking several layers of bi-directional (0/90) Nicalon 8-harness satin-weave fabric (Nippon Carbon, Tokyo, Japan). The average ®ber radius, as determined by quantitative microscopy, was 7.2 mm. The ®bers were coated by chemical vapor deposition with a thin layer of BN (approximately 100 nm) and then with a thicker layer of SiC (3 mm) onto the BN. The alumina matrix was in®ltrated by using a direct metal oxidation process. The preform was brought into contact with molten aluminum in air at 1000C. The aluminum reacted with the oxygen to form a matrix of porous Al2O3 which grew into the preform. The residual aluminum was ®nally removed from the composite material using a proprietary technique. The volume fraction of ®bers in the composite was 37%, while the SiC ®ber coating occupied another 37%. The Al2O3 matrix comprised 18% of the composite volume and the rest (8%) was porosity. More details of the manufacturing process and the microstructure can be found elsewhere [14±16]. Tensile and notched-beam specimens were machined from the plates. The tensile specimens had a dog-bone shape and were designed according to the speci®cations given by French and Wiederhorn [17]. The central part of the specimen had uniform width of 4 mm and a length of 15 mm. Load was applied through two ceramic pins which were introduced into circular holes machined in the specimen heads. The pins were attached to two alumina rods connected, respectively to the actuator and to the load cell of the mechanical testing machine. Two bi-directional hinges were inserted in the load train to avoid bending stresses during the tests. The experimental set-up is shown in Fig. 1. The fracture tests were carried out by three-point bending of notched prismatic bars. The loading span was 50 mm and the bar depth 10 mm. A notch of around 2 mm in length and 150 mm in radius was introduced with a thin diamond wire. The specimen and the loading ®xture were placed in a furnace for the elevated temperature tests. (see Fig. 1). Two B-type thermocouples were attached to the specimen for an accurate control of the temperature. The heating rate was 12C per min and each specimen was held at the test temperature for 1 h prior to testing to ensure uniform distribution of the temperature. Both tensile and fracture tests were performed under stroke control at a crosshead speed of 50 mm/min. The load (P) and the crosshead displacement of the testing machine with respect to the frame (n) were monitored during the tests, the latter through a LVDT transducer outside the furnace. In addition, the distance between two small pins glued in the central region of the tensile specimens was measured through a laser extensometer. This extensometer was also used during the fracture tests to measure the crack mouth opening displacement (CMOD) obtained from the distance between two alumina pins glued symmetrically to the notch mouth. Once broken, the fracture surfaces were sputtered with Au±Pd for 3 min before being observed in the scanning electron microscope. In addition, the specimens were sliced far away from the fracture surface with a low speed diamond saw, and the longitudinal sections were polished successively on diamond cloths of 40, 9, 3 and 1 mm grain size and ®nally on alumina of 0.3 mm grain size. They were cleaned for 30 min by ultrasound in acetone to remove the alumina from polishing, and studied in the optical microscope. 3. Experimental results 3.1. Mechanical properties The initial composite response during the tensile tests was linear. It was followed by a pronounced knee in the Fig. 1. Experimental set-up for the tensile tests at elevated temperature. 1068 J.A. CelemõÂn, J. LLorca / Composites Science and Technology 60 (2000) 1067±1076
J.A. Celemin, J. LLorca/ Composites Science and Technology 60(2000)1067-1076 1069 Table l problems with the specimen alignment precluded the Composite, matrix and fiber elastic moduli accurate determination of the modulus at 800 and Temperature(C 1000C and these were estimated using reasonable assumptions as indicated in [18]. The average values of the measured and estimated elastic moduli are shown in ErGPa) 135 Four P-CMOD curves corresponding to fracture tests at 25, 800, 1000, and 1200oC are depicted in Fig. 2(b) stress/ strain curve, which marked the transition to the They also presented an initial linear zone, which was non-linear regime induced by the onset of multiple followed by a non-linear region prior to the maximum matrix cracking. The results for the matrix cracking load in the tests at 25C. The observation of the notch non-linmear (defined as the stress at the beginning of the tip region through a telescope at this stage showed the are plotted in Fig. 2(a)as a function of the test tem- matrix cracks. The non -linear region was less noticeable perature. In addition, the elastic modulus was deter- at 800C and has practically disappeared at 1000 and mined from the initial slope of the stress/strain curve in 1200 C. Fracture took place by the propagation of a the specimens tested at 25 and 1200C. Experimental single dominant crack at these temperatures. 250 Matrix cracking stress Temperature 200 1000c 150 1200c 600H 200 200400600800100012001400 MOd (ut · Fracture energy 25 25 200400600800100012001400 200400600800100012001400 Temperature(.) strength,Ou;(b)load/crack mouth opening displacement curves; (c)nominal fracture toughness, Ko: (d)fracture energy,c, Nie"and hanical properties of the Al2O/Nicalon composite as a function of the test temperature: (a) tensile matrix crad
stress/strain curve, which marked the transition to the non-linear regime induced by the onset of multiple matrix cracking. The results for the matrix cracking stress, smc, (de®ned as the stress at the beginning of the non-linear regime) as well as for the tensile strength, su, are plotted in Fig. 2(a) as a function of the test temperature. In addition, the elastic modulus was determined from the initial slope of the stress/strain curve in the specimens tested at 25 and 1200C. Experimental problems with the specimen alignment precluded the accurate determination of the modulus at 800 and 1000C and these were estimated using reasonable assumptions as indicated in [18]. The average values of the measured and estimated elastic moduli are shown in Table 1. Four P-CMOD curves corresponding to fracture tests at 25, 800, 1000, and 1200C are depicted in Fig. 2(b). They also presented an initial linear zone, which was followed by a non-linear region prior to the maximum load in the tests at 25C. The observation of the notch tip region through a telescope at this stage showed the formation of a diuse damage zone, containing multiple matrix cracks. The non-linear region was less noticeable at 800C and has practically disappeared at 1000 and 1200C. Fracture took place by the propagation of a single dominant crack at these temperatures. Table 1 Composite, matrix and ®ber elastic moduli Temperature (C) 25 800 1000 1200 E (GPa) 73 69 52 45 Em(GPa) 145 137 98 80 Ef(GPa) 180 170 135 125 Fig. 2. Mechanical properties of the Al2O3/Nicalon composite as a function of the test temperature: (a) tensile matrix cracking stress, mc, and strength, u; (b) load/crack mouth opening displacement curves; (c) nominal fracture toughness, KQ; (d) fracture energy, GF. J.A. CelemõÂn, J. LLorca / Composites Science and Technology 60 (2000) 1067±1076 1069
J.A. Celemin, J. LLorca/ Composites Science and Technology 60(2000)1067-1076 The nominal fracture toughness, Ko, was obtained 25C [Fig 3(a)]. This indicates that the matrix cracks from the maximum load, Pu, and the initial notch did not propagate through the fibers, which remained gth, ao, according to [19 intact behind the crack tip bridging the crack surfaces They were broken in tension within the matrix as the KQ=5PuLa/BW/Y(a) (1 separation between the crack surfaces increased, and had to be pulled out prior to the complete specimen fracture. Fiber pull-out was also dominant in the speci where L is the span, B and w stand for the specimen mens tested at 800C [Fig 3(b)] and 1200oC [Fig 3(c)]. thickness and depth, respectively, a=ao/w, and Y(o)is although it should be noted that fibers broken in the a non-dimensional function given matrix crack plane were also observed in these cases Neither the fraction of pulled-out fibers nor the average 19179-12795a+3.3532a2-3.2260a3+1.2235a4 pull-out length were measured, but careful analyses of (1-a)3/(+2a) (2) This expression is valid for 0 <a<I when L/w=5. It should be noticed at this point that Ko depends on the specimen geometry and size because the length of the racture process zone around the crack tip was compar able to the specimen characteristic dimensions [20] Thus, Ko cannot be considered, strictly speaking,a material property but it is still a good parameter to estimate the influence of the temperature on the com posite toughness. This can also be analyzed through the fracture energy, GF, which stands for the energy spent 200pm to create a unit area of free surface. Assuming that the specimen fracture took place by the propagation of (b) crack from the notch to the back of the specimen, GE B(W-0Pdv where integral represents the area under the P-v The experimental results for the nominal fracture toughness and the fracture energy are plotted in Fig 2(c)and(d), respectively, as a function of the test tem- perature. Two main conclusions may be drawn from the results in Fig. 2. Firstly, the composite presented a duc tile behavior in the whole temperature range analyzed the matrix cracking stress was always significantly lower than the tensile strength, and the fracture toughness and the fracture energy were well above those measured in monolithic ceramics. Secondly, the strength and the toughness of the composite decreased significantly from ambient to 800oC and remained practically constant above this temperature 3. 2. Fractography. The analyses of the fracture surfaces in the scanning electron microscope corroborated the presence of a Fig 3. Fracture surface of the specimens tested at different tempera- weak fiber/matrix interface. Fibers protruding from the tures, showing fibers pulled out from the matrix: (a)25%C; (b)800oC fracture surface were observed in the specimens tested at (c)1200C
The nominal fracture toughness, KQ, was obtained from the maximum load, Pu, and the initial notch length, a0, according to [19] KQ 3 2 PuL 1=2 =BW3=2 Y 1 where L is the span, B and W stand for the specimen thickness and depth, respectively, a0=W, and Y( ) is a non-dimensional function given by Y 1:9179ÿ1:2795 3:3532 2ÿ3:2260 3 1:2235 4 1ÿ 3=2 12 2 This expression is valid for 0 < <1 when L=W 5. It should be noticed at this point that KQ depends on the specimen geometry and size because the length of the fracture process zone around the crack tip was comparable to the specimen characteristic dimensions [20]. Thus, KQ cannot be considered, strictly speaking, a material property but it is still a good parameter to estimate the in¯uence of the temperature on the composite toughness. This can also be analyzed through the fracture energy, GF, which stands for the energy spent to create a unit area of free surface. Assuming that the specimen fracture took place by the propagation of a crack from the notch to the back of the specimen, GF can be computed as GF 1 B W ÿ a0 Pdv 3 where the integral represents the area under the P±v curve. The experimental results for the nominal fracture toughness and the fracture energy are plotted in Fig. 2(c) and (d), respectively, as a function of the test temperature. Two main conclusions may be drawn from the results in Fig. 2. Firstly, the composite presented a ductile behavior in the whole temperature range analyzed: the matrix cracking stress was always signi®cantly lower than the tensile strength, and the fracture toughness and the fracture energy were well above those measured in monolithic ceramics. Secondly, the strength and the toughness of the composite decreased signi®cantly from ambient to 800C and remained practically constant above this temperature. 3.2. Fractography. The analyses of the fracture surfaces in the scanning electron microscope corroborated the presence of a weak ®ber/matrix interface. Fibers protruding from the fracture surface were observed in the specimens tested at 25C [Fig. 3(a)]. This indicates that the matrix cracks did not propagate through the ®bers, which remained intact behind the crack tip bridging the crack surfaces. They were broken in tension within the matrix as the separation between the crack surfaces increased, and had to be pulled out prior to the complete specimen fracture. Fiber pull-out was also dominant in the specimens tested at 800C [Fig. 3(b)] and 1200C [Fig. 3(c)], although it should be noted that ®bers broken in the matrix crack plane were also observed in these cases. Neither the fraction of pulled-out ®bers nor the average pull-out length were measured, but careful analyses of Fig. 3. Fracture surface of the specimens tested at dierent temperatures, showing ®bers pulled out from the matrix: (a) 25C; (b) 800C; (c) 1200C. 1070 J.A. CelemõÂn, J. LLorca / Composites Science and Technology 60 (2000) 1067±1076
J.A. Celemin, J. LLorca/Composites Science and Technology 60(2000)1067-1076 Table 2 Fraction of pulled-out fibers with different morphology on the fracture Temperature(°C)25 1000 ype I(%) Type Il(%) pe Ill (% These results can be used to compute the fiber failure probability, F, as a function of the fiber strength for Mirror each temperature. The strength data computed from the ist mirror radius were arranged in ascending order and 5 um corresponding failure probability, F=(i-0.5)/N, was assigned to each strength based on rank statistics, where Fig 4. Detail of the fracture surface of a pulled-out Nicalon SiC fiber, i is the rank and N is the total number of experimental showing the mirror-mist-hackle morphology data. The lowest strength was attributed to the fibers with a specular fracture surface(type Ill), following the analysis of Eckel and Bradt [22]. The highest strength the fracture surfaces led to two qualitative-but was assigned to the fibers with very short mirror radius unambiguous-conclusions. Firstly, the average fiber (type I). The results are shown in Fig. 5, and they are 1200C was shorter than in those tested at 25C pull-out length in the specimens tested at 1000 well approximated by a Weibull function of the form ondly, the fraction of fibers broken in the matrix crad plane was maximum at 800oC 3.3. In situ fiber strengi (5)where the parameter o* can be obtained from Fig. 5 If observed at higher magnification, the fracture sur- as the stress which gives a fracture probability equal to face of most of the pulled-out fibers showed the mor- 63%, and m* can be computed by the least squares fit phology depicted in Fig. 4, which includes a specular ting of (5) to the experimental results plotted in Fig. 5 circular region(mirror), an intermediate rough surface The values of os and m derived from the fracture mir (mist), and an abrupt external area(hackle). The mirror- ror data are not, in general, identical to the true in situ misf-hackle topography is typical on the fracture sur- fiber characteristic strength, ac, and Weibull modulus m, face of Nicalon Sic fibers broken in tension. An empirical relationship between the fracture mirror dius, am(expressed in m), and the fiber strength, s(in 100 MPa), has been shown by several authors to be of the 2.51 for Nicalon SiC fibers [21-23]. to determine the in situ fiber strength, the fracture surface of approximately 130 pulled-out fibers was analyzed for each temperature The majority of the fibers exhibited a distinct mirror- 25 misf-hackle structure(type II), where the mirror radius could be easily determined. However, this radius was too short to be accurately measured in a small fraction 1200c of fibers (type I) while no distinct fracture mirror boundary was seen in other fibers and the whole fiber 1000 2000 fracture surface was specular(type IIl). The proportion Fiber Strength, S(MPa) of fibers with type I, II and Ill fracture surfaces is Fig. 5. Cumulative fracture probability of the fibers, F, at different shown in Table 2 as a function of the test temperature temperatures as obtained from the fracture mirror data
the fracture surfaces led to two qualitative±±but unambiguous Ð conclusions. Firstly, the average ®ber pull-out length in the specimens tested at 1000 and 1200C was shorter than in those tested at 25C. Secondly, the fraction of ®bers broken in the matrix crack plane was maximum at 800C. 3.3. In situ ®ber strength If observed at higher magni®cation, the fracture surface of most of the pulled-out ®bers showed the morphology depicted in Fig. 4, which includes a specular circular region (mirror), an intermediate rough surface (mist), and an abrupt external area (hackle). The mirror± mist±hackle topography is typical on the fracture surface of Nicalon SiC ®bers broken in tension. An empirical relationship between the fracture mirror radius, am (expressed in m), and the ®ber strength, S (in MPa), has been shown by several authors to be of the form, S 2:51 am p 4 for Nicalon SiC ®bers [21±23]. To determine the in situ ®ber strength, the fracture surface of approximately 130 pulled-out ®bers was analyzed for each temperature. The majority of the ®bers exhibited a distinct mirror± mist±hackle structure (type II), where the mirror radius could be easily determined. However, this radius was too short to be accurately measured in a small fraction of ®bers (type I) while no distinct fracture mirror boundary was seen in other ®bers and the whole ®ber fracture surface was specular (type III). The proportion of ®bers with type I, II and III fracture surfaces is shown in Table 2 as a function of the test temperature. These results can be used to compute the ®ber failure probability, F, as a function of the ®ber strength for each temperature. The strength data computed from the mirror radius were arranged in ascending order and a corresponding failure probability, F=(iÿ0.5)/N, was assigned to each strength based on rank statistics, where i is the rank and N is the total number of experimental data. The lowest strength was attributed to the ®bers with a specular fracture surface (type III), following the analysis of Eckel and Bradt [22]. The highest strength was assigned to the ®bers with very short mirror radius (type I). The results are shown in Fig. 5, and they are well approximated by a Weibull function of the form F 1 ÿ exp ÿ S c " # m 5 (5) where the parameter c can be obtained from Fig. 5 as the stress which gives a fracture probability equal to 63%, and m can be computed by the least squares ®tting of (5) to the experimental results plotted in Fig. 5. The values of c and m derived from the fracture mirror data are not, in general, identical to the true in situ ®ber characteristic strength, c, and Weibull modulus m, Table 2 Fraction of pulled-out ®bers with dierent morphology on the fracture surface Temperature (C) 25 800 1000 1200 Type I (%) 22 20 24 31 Type II (%) 66 63 69 58 Type III (%) 12 17 7 11 Fig. 4. Detail of the fracture surface of a pulled-out Nicalon SiC ®ber, showing the mirror±mist±hackle morphology. Fig. 5. Cumulative fracture probability of the ®bers, F, at dierent temperatures as obtained from the fracture mirror data. J.A. CelemõÂn, J. LLorca / Composites Science and Technology 60 (2000) 1067±1076 1071
J.A. Celemin, J. LLorca/ Composites Science and Technology 60(2000)1067-1076 Table 3 to the classical analysis of Aveston et al. [24], matrix Weibull parameters for the in situ fiber strength cracking in the 0o plies develops as a consequence of Temperature(C) 25 load transfer from the fibers when the composite stress attains the matrix cracking stress, ome, which is given 1/ which characterize the fiber-strength distribution me according to the Weibull model. They have to be cor- rected to account for screening of flaws over a finite fiber length on either side of a fiber-failure site. Follow- where ym stands for the matrix fracture energy, Em and ng the procedure developed by Curtin [23), the cor- Er represent the matrix and fiber elastic moduli, respec. rected values of o and m were estimated and a tively, and R is the fiber radius. This expression was presented in Table 3 for each temperature. The fiber originally obtained for uniaxially-reinforced compo Weibull modulus at 25C was very low( 2), indicating sites, and its extension to bi-directional fabrics is that the fiber defect distribution in the as-received achieved by using only the volume fraction of fibers that material was very wide and remained so at elevated are oriented in the loading direction(18.5%)in(6) temperature. The characteristic strength remained con- As multiple matrix cracking develops, the slip zones stant up to 800c but decreased rapidly above this from neighboring cracks overlap and produce a shield- temperature. ing effect. When shielding proceeds to completion, the average spacing of the matrix cracks, dme, is expressed 3.4. Interfacial sliding resistance specimens in the optical microscope showed the pre- dme=1. 6R(-'ymErE-7/3 Observation of the longitudinal sections of the tensile fter sence of arrays of parallel matrix cracks oriented per pendicularly to the loading axis(Fig. 6). These matrix cracks were clearly defined in matrix regions within the and the interfacial sliding resistance, L, is then obtained fiber plies that were parallel to the loading axis(0 from(6) and(7)as plies), and became less regular as they interacted with fiber plies perpendicular to the loading axis(90 plies) The average matrix crack spacing, dmc, in the 0 fiber r=1.761 plies was computed from the distance between adjacent cracks measured on more than 30 crack pairs for each fump erature. The results are shown in Table 4 as a where it is only necessary to determine the matrix elastic tion of the test temperature modulus as a function of temperature. This was esti The average crack spacing is a useful parameter to mated from the composite and fiber moduli assumin estimate the interfacial sliding resistance, T. According that the composite modulus was dominated by the 0o plies and neglecting the contribution of the 90 plies This latter hypothesis was supported by the presence of cracks in the matrix prior to testing due to the differ ences in the thermal expansion coefficients of Al2O3 and Nicalon SiC fibers [15, 16]. As the volume of the SiC fiber coating was considerable, the matrix was taken as the sum of the Al,O3 matrix(18%)and the Sic coating (37%), which gave a total matrix volume fraction fm=0.55. Assuming that the isostrain model is ade quate to simulate the elastic deformation within the 0o plies, the composite modulus is given by Table 4 Average matrix crack spacing, dmc, and interfacial sliding resistance,r Fig. 6. Longitudinal section of the composite tested at 1200C show dne(pm)(±SD.)215±61 9±57160±43 ng the presence of arrays of parallel matrix cracks. Loading axis was I(MPa)
which characterize the ®ber-strength distribution according to the Weibull model. They have to be corrected to account for screening of ¯aws over a ®nite ®ber length on either side of a ®ber-failure site. Following the procedure developed by Curtin [23], the corrected values of c and m were estimated and are presented in Table 3 for each temperature. The ®ber Weibull modulus at 25C was very low (2), indicating that the ®ber defect distribution in the as-received material was very wide and remained so at elevated temperature. The characteristic strength remained constant up to 800C but decreased rapidly above this temperature. 3.4. Interfacial sliding resistance Observation of the longitudinal sections of the tensile specimens in the optical microscope showed the presence of arrays of parallel matrix cracks oriented perpendicularly to the loading axis (Fig. 6). These matrix cracks were clearly de®ned in matrix regions within the ®ber plies that were parallel to the loading axis (0 plies), and became less regular as they interacted with ®ber plies perpendicular to the loading axis (90 plies). The average matrix crack spacing, dmc, in the 0 ®ber plies was computed from the distance between adjacent cracks measured on more than 30 crack pairs for each temperature. The results are shown in Table 4 as a function of the test temperature. The average crack spacing is a useful parameter to estimate the interfacial sliding resistance, . According to the classical analysis of Aveston et al. [24], matrix cracking in the 0 plies develops as a consequence of load transfer from the ®bers when the composite stress attains the matrix cracking stress, mc, which is given by: mc E 6 mEff 2 EE2 mR 1 ÿ f 1=2 6 where m stands for the matrix fracture energy, Em and Ef represent the matrix and ®ber elastic moduli, respectively, and R is the ®ber radius. This expression was originally obtained for uniaxially-reinforced composites, and its extension to bi-directional fabrics is achieved by using only the volume fraction of ®bers that are oriented in the loading direction (18.5%) in (6). As multiple matrix cracking develops, the slip zones from neighboring cracks overlap and produce a shielding eect. When shielding proceeds to completion, the average spacing of the matrix cracks, dmc, is expressed by [25]: dmc 1:6R 1 ÿ f 2 mEfEm f 2ER 1=3 7 and the interfacial sliding resistance, t, is then obtained from (6) and (7) as, 1:761 1 ÿ f f R 2 Em E mc dmc 8 where it is only necessary to determine the matrix elastic modulus as a function of temperature. This was estimated from the composite and ®ber moduli assuming that the composite modulus was dominated by the 0 plies and neglecting the contribution of the 90 plies. This latter hypothesis was supported by the presence of cracks in the matrix prior to testing due to the dierences in the thermal expansion coecients of Al2O3 and Nicalon SiC ®bers [15,16]. As the volume of the SiC ®ber coating was considerable, the matrix was taken as the sum of the Al2O3 matrix (18%) and the SiC coating (37%), which gave a total matrix volume fraction fm 0:55. Assuming that the isostrain model is adequate to simulate the elastic deformation within the 0 plies, the composite modulus is given by Table 3 Weibull parameters for the in situ ®ber strength Temperature (C) 25 800 1000 1200 c(GPa) 1.82 1.78 1.44 1.36 m 2.0 2.3 2.9 2.6 Fig. 6. Longitudinal section of the composite tested at 1200C showing the presence of arrays of parallel matrix cracks. Loading axis was vertical. Table 4 Average matrix crack spacing, dmc, and interfacial sliding resistance, Temperature (C) 25 800 1000 1200 dmc (mm) ( S.D.) 21561 16651 18957 16043 (MPa) 23.8 23.0 17.9 17.5 1072 J.A. CelemõÂn, J. LLorca / Composites Science and Technology 60 (2000) 1067±1076
(9) Curtin [27] to predict the tensile strength and tough J.A. Celemin, J. LLorca/ Composites Science and Technology 60(2000)1067-1076 E=2LEd+Em/mI of FRC with a weak fiber /matrix interface. The model, which has been extensively validated [25, 28], assumes where f is the fraction of fibers(37%)and the 1/2 factor that fiber failure is well represented by the Weibull sta comes from neglecting the contribution of the 90 plies. tistics and that the fibers fracture independently during The elastic moduli of the Nicalon sic fiber as a function deformation. In addition the weak interface relieves the of temperature can be found in [26]. They are shown in stress concentration at the fiber breaks, and thus global Table 1, where the corresponding matrix moduli com- load redistribution takes place when fiber fracture compute the interfacial sliding resistance through(8), strength, Ou, is given by iOns, the composite tensile puted from(9)can be found as well. They were used to occurs. On these assumpti which is shown in Table 4 as a function of temperature ambient temperature computed from the average crack u=foc m+2 spacing was in good agreement with the results obtained by Heredia et al. [15] on this material where f is the fraction of fibers oriented in the loading axis(0. 175 in our case) and e is the characteristic fiber strength as obtained from the fracture mirror analysis. Curtin also computed the fracture energy as the sum of the energy spent in breaking the fibers(Gb)plus the The mechanical tests showed a marked reduction in energy necessary to pull out the broken fibers from the the composite strength and toughness at intermediate matrix(Gp), both per unit of free surface Mathemati- and elevated temperatures, albeit the material still lly [27] exhibited a non-linear behaviour. In agreement with these results, the fractographic observations found fiber GF=Gb+G=0.4 RoaEm(-/+dmoR(D) pull-out in the fracture surfaces in the whole tempera ire range although the fraction of fibers broken in the matrix crack plane was maximum at 800oC. Finally, The total fra acture ener gy is dominated by the pull-out neither the fiber strength nor the interfacial sliding contribution, which is usually one order of magnitude resistance varied significantly between 25 and 800C but higher than Gb. The function (m) can be found in the both decreased above this temperature. These observa- cited reference. Its value depends on whether the speci s can be used to understand the embrittlement men failure took place by the propagation of one mechanisms at different temperatures with the help of dominant matrix crack from the notch tip to the back of the successful micromechanical model developed by the specimen, or was better represented by a diffuse damage zone containing multiple matrix cracks, which 250 The model results for the tensile strength are com- pared with the experimental values in Fig. 7. Regardless of the experimental scatter, the model predictions are in 200 reasonably good agreement with the experiments at ambient and elevated temperatures(1000 and 1200oC) This indicates that the main model hypotheses were 150 fulfilled at these temperatures: the fibers were fractured randomly along the specimen and the load released by broken fibers was homogeneously redistributed among the intact fibers. The reduction of the composite Temperature strength at elevated temperature was caused by the 20- 25% reduction of the in situ fiber strength, as compared with the fibers in the as-received composite. The strength degradation of the Nicalon SiC fibers at 口1000gc 1000oC and above is a well documented phenomenon [10, 26, 29-31], which is mainly produced by the nuclea 50100150 tion and growth of defects in the fiber surface. g (MPa), experimental According to the model, the specimen tested at 800C should exhibit a tensile strength very close to that mea Fig. 7. Model predictions and experimental results for the composite sured at ambient temperature. This was not so, however and the composite strength was significantly lower
E 1 2 Eff Emfm 9 where f is the fraction of ®bers (37%) and the 1/2 factor comes from neglecting the contribution of the 90 plies. The elastic moduli of the Nicalon SiC ®ber as a function of temperature can be found in [26]. They are shown in Table 1, where the corresponding matrix moduli computed from (9) can be found as well. They were used to compute the interfacial sliding resistance through (8), which is shown in Table 4 as a function of temperature. It is worth noting that interfacial sliding resistance at ambient temperature computed from the average crack spacing was in good agreement with the results obtained by Heredia et al. [15] on this material. 4. Discussion The mechanical tests showed a marked reduction in the composite strength and toughness at intermediate and elevated temperatures, albeit the material still exhibited a non-linear behaviour. In agreement with these results, the fractographic observations found ®ber pull-out in the fracture surfaces in the whole temperature range, although the fraction of ®bers broken in the matrix crack plane was maximum at 800C. Finally, neither the ®ber strength nor the interfacial sliding resistance varied signi®cantly between 25 and 800C but both decreased above this temperature. These observations can be used to understand the embrittlement mechanisms at dierent temperatures with the help of the successful micromechanical model developed by Curtin [27] to predict the tensile strength and toughness of FRC with a weak ®ber/matrix interface. The model, which has been extensively validated [25,28], assumes that ®ber failure is well represented by the Weibull statistics and that the ®bers fracture independently during deformation. In addition, the weak interface relieves the stress concentration at the ®ber breaks, and thus global load redistribution takes place when ®ber fracture occurs. On these assumptions, the composite tensile strength, u, is given by [27], u fc 2 m 2 1= m1 m 1 m 2 10 where f is the fraction of ®bers oriented in the loading axis (0.175 in our case) and c is the characteristic ®ber strength as obtained from the fracture mirror analysis. Curtin also computed the fracture energy as the sum of the energy spent in breaking the ®bers (Gb) plus the energy necessary to pull out the broken ®bers from the matrix (Gp), both per unit of free surface. Mathematically [27], GF Gb Gp 0:4f R 3 c Em 1 ÿ f EEf f ' m 12 2 c R 11 The total fracture energy is dominated by the pull-out contribution, which is usually one order of magnitude higher than Gb. The function ' m can be found in the cited reference. Its value depends on whether the specimen failure took place by the propagation of one dominant matrix crack from the notch tip to the back of the specimen, or was better represented by a diuse damage zone containing multiple matrix cracks, which grew from the notch tip. The model results for the tensile strength are compared with the experimental values in Fig. 7. Regardless of the experimental scatter, the model predictions are in reasonably good agreement with the experiments at ambient and elevated temperatures (1000 and 1200C). This indicates that the main model hypotheses were ful®lled at these temperatures: the ®bers were fractured randomly along the specimen and the load released by broken ®bers was homogeneously redistributed among the intact ®bers. The reduction of the composite strength at elevated temperature was caused by the 20± 25% reduction of the in situ ®ber strength, as compared with the ®bers in the as-received composite. The strength degradation of the Nicalon SiC ®bers at 1000C and above is a well documented phenomenon [10,26,29±31], which is mainly produced by the nucleation and growth of defects in the ®ber surface. According to the model, the specimen tested at 800C should exhibit a tensile strength very close to that measured at ambient temperature. This was not so, however and the composite strength was signi®cantly lower Fig. 7. Model predictions and experimental results for the composite tensile strength. J.A. CelemõÂn, J. LLorca / Composites Science and Technology 60 (2000) 1067±1076 1073
J.A. Celemin, J. LLorca/ Composites Science and Technology 60(2000)1067-1076 although the fibers did not exhibit any degradation. The different temperatures confirmed that distributed crack- examination of the fracture surfaces [see Fig. 3(b)] ing was prevalent at 25C and that failure took place by showed that the fraction of fibers broken in the matrix the propagation of a single crack at 800oC and above. lane was maximum at this temperature, which may This change in the fracture mechanism between ambient indicate that the ingress of oxygen in the composite led and elevated temperatures was also reported by gomina to the oxidation of the fiber/matrix interface in a sig- et al. [32] and Nair and Wang [33] in 2-D woven FRC, nificant number of fibers. The matrix cracks did propa- and it stands for the main cause of the reduction in the gate through these fibers, and this localized the damage composite toughness at intermediate and elevated tem in a given section of the composite, promoting the early perature. Damage is accumulated very quickly at the specimen rupture. Although this behaviour may be sur- notch tip during the fracture tests, and this facilitates prising, two previous investigations on Nicalon-fiber the access of oxygen to the interior of the sample, loca- composites coated with Bn had similar results. Sun et lizing the fracture above and below one single dominant al. [12] attributed the interface degradation at 600 C to crack. It should be noted that the change in the fracture the fact that the Nicalon oxidation rate at this tem- mode between 25 and 800C was responsible for the perature was not enough to seal the stress-induced marked decrease in the composite fracture energy even cracks, and nothing impeded the access of oxygen to the though the relevant microstructural parameters(fiber nterior of the composite. In contrast, the rapid silicate strength and interfacial sliding resistance) were practi formation at 950C sealed the cracks and protected the cally constant in this temperature range interfaces within the composite. Morscher [30 also observed a mild embrittlement of the composites in the temperature range 900-1 100C, which was ascribed to 5. Conclusions the fusion of the fiber to the SiC matrix by glass forme in the interphase region. This process was not observed The strength and toughness of a 2-D woven nicalon however,at1200°C. Al2O3-matrix composite were measured at ambient The model results for the fracture energy are plotted intermediate(800C), and elevated(1000-1200oC)tem in Fig 8, together with the experimental data. The pre- peratures. The composite exhibited a non-linear beha dictions in Fig 8(a) stand for multiple matrix cracking, viour in the whole temperature range but the whereas those in Fig 8(b) represent the response when mechanical properties were significantly degraded at fracture is due to a single matrix crack. The ambient 800 C and above. Quantitative microscopy techniques model predictions in the former case, while those at facial sliding resistance as a function of temperature. intermediate and elevated temperatures followed the Both parameters remained practically constant between behaviour for a single matrix crack. The in situ obser- 25 and 800C but showed a noticeable reduction above vation of the damage zone during the fracture tests at this temperature Multiple Matrix Cracking Single Matrix Cracking Temperature 25c B 800%c 30 口10009c 1200c 20 5 10009c 1200c (b) 0510152025303540 0510152025303540 G.(kJ/ms) tal G(kJ/m2),experimental Fig 8. Model predictions and experimental results for the composite fracture energy: (a)multiple matrix cracking; (b) single matrix cracking
although the ®bers did not exhibit any degradation. The examination of the fracture surfaces [see Fig. 3(b)] showed that the fraction of ®bers broken in the matrix plane was maximum at this temperature, which may indicate that the ingress of oxygen in the composite led to the oxidation of the ®ber/matrix interface in a signi®cant number of ®bers. The matrix cracks did propagate through these ®bers, and this localized the damage in a given section of the composite, promoting the early specimen rupture. Although this behaviour may be surprising, two previous investigations on Nicalon±®ber composites coated with BN had similar results. Sun et al. [12] attributed the interface degradation at 600C to the fact that the Nicalon oxidation rate at this temperature was not enough to seal the stress-induced cracks, and nothing impeded the access of oxygen to the interior of the composite. In contrast, the rapid silicate formation at 950C sealed the cracks and protected the interfaces within the composite. Morscher [30] also observed a mild embrittlement of the composites in the temperature range 900±1100C, which was ascribed to the fusion of the ®ber to the SiC matrix by glass formed in the interphase region. This process was not observed, however, at 1200C. The model results for the fracture energy are plotted in Fig. 8, together with the experimental data. The predictions in Fig. 8(a) stand for multiple matrix cracking, whereas those in Fig. 8(b) represent the response when fracture is due to a single matrix crack. The ambient temperature results were in good agreement with the model predictions in the former case, while those at intermediate and elevated temperatures followed the behaviour for a single matrix crack. The in situ observation of the damage zone during the fracture tests at dierent temperatures con®rmed that distributed cracking was prevalent at 25C and that failure took place by the propagation of a single crack at 800C and above. This change in the fracture mechanism between ambient and elevated temperatures was also reported by Gomina et al. [32] and Nair and Wang [33] in 2-D woven FRC, and it stands for the main cause of the reduction in the composite toughness at intermediate and elevated temperature. Damage is accumulated very quickly at the notch tip during the fracture tests, and this facilitates the access of oxygen to the interior of the sample, localizing the fracture above and below one single dominant crack. It should be noted that the change in the fracture mode between 25 and 800C was responsible for the marked decrease in the composite fracture energy even though the relevant microstructural parameters (®ber strength and interfacial sliding resistance) were practically constant in this temperature range. 5. Conclusions The strength and toughness of a 2-D woven Nicalon/ Al2O3-matrix composite were measured at ambient, intermediate (800C), and elevated (1000±1200C) temperatures. The composite exhibited a non-linear behaviour in the whole temperature range but the mechanical properties were signi®cantly degraded at 800C and above. Quantitative microscopy techniques were used to estimate the ®ber strength and the interfacial sliding resistance as a function of temperature. Both parameters remained practically constant between 25 and 800C but showed a noticeable reduction above this temperature. Fig. 8. Model predictions and experimental results for the composite fracture energy: (a) multiple matrix cracking; (b) single matrix cracking. 1074 J.A. CelemõÂn, J. LLorca / Composites Science and Technology 60 (2000) 1067±1076
J.A. Celemin, J. LLorca/Composites Science and Technology 60(2000)1067-1076 The experimental results were compared with the erties of ceramic-matrix composites. predictions of a micromechanical model based on two JOM,May1993:57-63 assumptions: fibers fracture independently and their [7 Naslain R. Lamon J, Doumeingts D, editors. High temperatur load is redistributed uniformly among the intact fibers ceramic matrix composites. Bordeaux: Woodhead Publishing The model predictions, as well as the fractographic [8) Evans AG. Naslain R, editors. High temperature ceramic matrix observations, showed that the degradation mechanisms ceramic transactions. voL. 58. Westerville: The were different at intermediate and elevated tempera- American Ceramic Society tures. While the reduction in the composite strength at 9 Ihara K, Nakano K, Sekino T, Yasuda E, editors. Hig 1000 and 1200oC was caused mainly by the degradation perature ceramic matrix composites Ill. Uetikon-Zuerich: T Tech Publications. 1998 of the Nicalon fiber strength, embrittlement at 800C [10 LLorca J, Elices M, Celemin JA. Toughness and microstructural was due to the oxidation of the fiber/matrix interface in degradation at high temperature in SiC fiber-reinforced ceramics. a noticeable number of fibers. These were broken by the Acta mater1998:46:2441-53. propagation of the matrix cracks and did not contribute [ll] Heredia FE, McNulty JC, Zok FW, Evans AG. Oxidation to the composite strength Soc1995;78:2097-100 The composite fracture energy also decreased visibly [12] Sun EY, Lin HT, Brennan J Intermediate-temperature envir- at intermediate and elevated temperatures. The experi nmental effects on boron nitride- coated silicon carbide fiber. mental results were in agreement with the model pre- inforced glass-ceramic composites. J Am Ceram Soc dictions, given that fracture at 25C occurred by the 1997:80:609-14 propagation of a damage zone containing multiple [3 Pharaoh MW, Daniel AM. Lewis MH. Stability of interfaces in calcium aluminosilicate matrix/Nicalon SiC fibre composites cracks from the notch tip, while high temperature fail- Mater Sci Let 1993: 12: 998-1001 ure was localized around one single crack. These differ [14] Fareed AS, Schiroky GH. Kennedy CR. Development of BN/SiC ences were confirmed by the in situ observation of duplex fiber coatings for fiber reinforced alumina matrix compo- damage progress during the tests. The rapid damage sites fabricated by directed metal oxidation. Ceram Eng Sci Proc 1993:14:794801 accumulation at the notch tip favoured the access of [15] Heredia FE, Evans AG, Anderson CA. Tensile and shear prop- oxygen to the interior of the sample, and this led to erties of continuous fiber-reinforced SiC/Al2O3 composites pro- interface oxidation at intermediate and elevated tem- cessed by melt oxidation. J Am Ceram Soc 1995: 78: 2790-800 peratures, localizing the fracture above and below one [16] Celemin JA, Pastor JY, LLorca J,Elices M. Mechanical be single dominant crack. These changes in the failure at 20 and 1200.C of Nicalon- silicon- carbide-fiber-reinforced mode were responsible for the reduction in the compo- osites. J Am Ceram Soc 1997- 80: 2569-80 [17 French JD, Wiederhorn SM. Tensile specimens from ceramic site toughness at 800C and above. mponents. J Am Ceram Soc 1996: 79: 550-2 [18 CelemIn JA. The relationship between microstructure and echanical properties at high temperature in Sic-fiber reinforced Acknowledgements Al2O3. Doctoral thesis, Polytechnic University of Madrid, 1998 [19 Guinea GV, Pastor JY, Planas J, Elices M. Stress intensity factor compliance, and CMOD for a general three-point bend beam. The authors are indebted to dr. JY. Pastor for his J Fract1998;89:103-16 help in the mechanical tests at elevated temperature. [20] LLorca J, Elices M. Influence of specimen geometry and size on his investigation was supported by CICYT(Spain he fracture of fiber-reinforced ceramic composites. Engng Fract and NATO through grants MAT 95-787 and CRG Mech1993;44:341-58 941033, respectively 21] Sawyer LC, Jamieson M, Brikowsky D, Haider MI, Chen Rt. rength, structure and fracture properties of ceramic fibers duced from polymeric precursors: I, base-line studies. JAm Ceram Soc1987;70:798-810 References 22] Eckel AJ. Bradt RC Strength distribution of reinforcing fibers in a Nicalon fiber/chemically vapor infiltrated silicon carbide matrix composite J Amer Ceram Soc 1989: 72: 455-8 [] Mazdiyasni ks, editor. Fiber-reinforced ceramic composites. [23) Curtin WA In situ fiber strengths in ceramic-matrix composites Park Ridge (NJ): Noyes Publications, 1990. from fracture mirrors. J Am Ceram Soc 1994: 77: 1075-8 [2] Nair SV, Jakus K, editors. High temperature mechanical beha- [24] Aveston J. Cooper GA, Kelly A. Single and multiple fracture, the properties of fiber composites. IPC Science and Technology Press, Teddington (UK): National Physical Laboratory 3] Lehman RL, El-Rahaiby SK, Wachtman JB, editors. Handbook [25] Evans AG, Domergue JM, Vagaggini E. J Am Ceram Soc 1994:77:1425-35 berville(OH): American Ceramic Society, 1995. 6 Physer DJ, Coretta KC, Hodder RS, Tressler RE. J Am Ceram [4 Prewo KM, Johnson B, Starret S Silicon carbide fibre-reinforced Soc1989;72:284-8 glass-ceramic composite tensile behaviour at elevated tempera. [27 Curtin WA. Theory of mechanical properties of ceramic-matrix J Mater Sci 1989: 24: 1373-9 omposites. J Am Ceram Soc 1991: 74: 2837-45 5 Gomina M, Fourvel P, Rouillon MH. High tempera 28 Stawovy RH, Kampe SL, Curtin WA. Mechanical behavior of mechanical behaviour of an uncoated Sic-Sic composite and blackglas ceramic matrix composites. Acta Mater rial. J Mater Sci 1991- 26: 1891-6 [6]Woodford DA, Van Steele DR, Brehm JA, Timms LA, Palko JE, [29] Berger MH, Hochet N, Bunsell AR Microstructure and thermo-
The experimental results were compared with the predictions of a micromechanical model based on two assumptions: ®bers fracture independently and their load is redistributed uniformly among the intact ®bers. The model predictions, as well as the fractographic observations, showed that the degradation mechanisms were dierent at intermediate and elevated temperatures. While the reduction in the composite strength at 1000 and 1200C was caused mainly by the degradation of the Nicalon ®ber strength, embrittlement at 800C was due to the oxidation of the ®ber/matrix interface in a noticeable number of ®bers. These were broken by the propagation of the matrix cracks and did not contribute to the composite strength. The composite fracture energy also decreased visibly at intermediate and elevated temperatures. The experimental results were in agreement with the model predictions, given that fracture at 25C occurred by the propagation of a damage zone containing multiple cracks from the notch tip, while high temperature failure was localized around one single crack. These dierences were con®rmed by the in situ observation of damage progress during the tests. The rapid damage accumulation at the notch tip favoured the access of oxygen to the interior of the sample, and this led to interface oxidation at intermediate and elevated temperatures, localizing the fracture above and below one single dominant crack. These changes in the failure mode were responsible for the reduction in the composite toughness at 800C and above. Acknowledgements The authors are indebted to Dr. J.Y. Pastor for his help in the mechanical tests at elevated temperature. This investigation was supported by CICYT (Spain), and NATO through grants MAT 95-787 and CRG- 941033, respectively. 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