Acta mater. Vol. 46. No. 7, pp. 2441-2453, 1998 blished by Elsevier Science Ltd. All rights reserved PI:Sl359-645497)00402 359645498s1900+0.00 TOUGHNESS AND MICROSTRUCTURAL DEGRADATION AT HIGH TEMPERATURE IN SIC FIBER-REINFORCED CERAMICS J LLORCA, M. ELICES and J. A. CELEMIN Department of Materials Science, Polytechnic University of Madrid, ETS de Ingenieros de Caminos. 28040 Madrid, spain Abstract- The fracture behavior of three different SiC fiber-reinforced ceramics at emperature in air is studied. The fracture properties were obtained by Flexure tests on notched heams and approach. The microstructi terfacial and fiher degradation, were analyzed in each material and related to the ms and to the degradation in toughness and fracture resistance. Finally, the prove the high temperature toughness of fiber-reinforced ceramics are briefly discussed 1 INTRODUCTION systems result in exceptionally high levels of tough mh ness and strength, their properties drop very quickl h combine the properti ramics with a above 800C in oxidizing environments due to damage tolerant, ductile behavior has been one of matrix softening and fiber/matrix reactions [5, 6] the most active research topics in materials science in the last two decades. The driving forces for these stable ceramic matrices, and several processing tech e the potential gains in efficien niques(chemical vapor infiltration direct-metal oxi- and in power output of thermal engines at higher dation, polymer pyrolysis, melt infiltration, etc. [] s,and the interest in tough amic ma- were developed to introduce the ceramic matrix terials increased as further improvements in the into the fiber preform These efforts were rewarded with a number of working temperature of Ni-based superalloys were fiber-reinforced ceramics(FRC) with excellent frac (around 1250C). In principle, ceramics are the ture resistance and non-linear stress-strain curve at ideal candidates to substitute Ni-based superalloys mation and fracture in these materials were an as high temperature structural materials. They exhi tail,and para bit very high melting points, excellent chemical micromechanical models which related the macro- stability and wear and creep resistance as well scopic behavior t he microstructural low density. Their main drawback in structural ap- parameters (8,9). However, most of this work was plications lies in their reduced ductility and fracture carried out in the ambient temperature regime, toughness, which makes Ic components prone where FRC are unlikely to be used. The amount of to catastrophic failure. The addition of fibers to cer- work on their elevated temperature performance is amics has been known for many years to be one still limited, and this is especially true for the frac- way of improving these properties. The develop- ture toughness and fracture resistance. Only a hand- ment of fiber- reinforced cements Instance, is ful of investigations on these topics are available in undoubtedly a good example of this approach the open literature [6, 10-16 It was demonstrated at the beginning of the 70s It should also be noted that the fracture beh hat high performance C bers could be successfully of FRC at elevated temperature can be considered incorporated into ceramic glasses to achieve high a litmus test because the presence of a pre-existing strength, tough composite materials [1-4]. These notch or crack facilitates the entry of oxygen into early developments were not carried further due to the composite and accelerates any degradation pro the oxidative instability of C fibers at high tempera- cess which might occur. In addition, it is necessary ture. As the new oxidation-resistant SiC bers de- to find whether the present materials are able to rived by pyrolysis of organometallic precursors achieve the minimum toughness values of became available at the end of this decade, a new 15 MPa. m at 1350C laid down by the industry to generation appeared of glass and glass-ceramic introduce these new materials as structural com- matrices reinforced with SiC fibers [5]. while these ponents in gas turbines [17]. The present knowledge
Pergamon Acia mater. Vol. 46, No. I, pp. 2441-2453, 1998 0 1998 Acta Metallurgica Inc. Published by Elsevier Science Ltd. All rights reserved Printed in Great Britain PII: S1359-6454(97)00402-3 1359-6454/98 $19.00 + 0.00 TOUGHNESS AND MICROSTRUCTURAL DEGRADATION AT HIGH TEMPERATURE IN SIC FIBER-REINFORCED CERAMICS J. LLORCA, M. ELICES and J. A. CELEMiN Department of Materials Science, Polytechnic University of Madrid, ETS de Ingenieros de Caminos, 28040 Madrid, Spain Abstract-The fracture behavior of three different Sic fiber-reinforced ceramics at ambient and elevated temperature in air is studied. The fracture properties were obtained by flexure tests on notched beams and the increase in fracture resistance with crack length was determined using an equivalent elastic crack approach. The microstructural changes during elevated temperature exposure, which could be divided into interfacial and fiber degradation, were analyzed in each material and related to the fracture micromechanisms and to the degradation in toughness and fracture resistance. Finally, the different strategies to improve the high temperature toughness of fiber-reinforced ceramics are briefly discussed. 0 1998 Acta Metallurgica Inc 1. INTRODUCTION The development of new ceramic-based materials which combine the properties of ceramics with a damage tolerant, ductile behavior has been one of the most active research topics in materials science in the last two decades. The driving forces for these investigations were the potential gains in efficiency and in power output of thermal engines at higher temperatures, and the interest in tough ceramic materials increased as further improvements in the working temperature of Ni-based superalloys were hindered by the proximity to their melting point (around 1250°C). In principle, ceramics are the ideal candidates to substitute Ni-based superalloys as high temperature structural materials. They exhibit very high melting points, excellent chemical stability and wear and creep resistance as well as low density. Their main drawback in structural applications lies in their reduced ductility and fracture toughness, which makes ceramic components prone to catastrophic failure. The addition of fibers to ceramics has been known for many years to be one way of improving these properties. The development of fiber-reinforced cements, for instance, is undoubtedly a good example of this approach. systems result in exceptionally high levels of toughness and strength, their properties drop very quickly above 800°C in oxidizing environments due to matrix softening and fiber/matrix reactions [5,6]. The emphasis was then focussed on using more stable ceramic matrices, and several processing techniques (chemical vapor infiltration, direct-metal oxidation, polymer pyrolysis, melt infiltration, etc. [7]) were developed to introduce the ceramic matrix into the fiber preform. These efforts were rewarded with a number of fiber-reinforced ceramics (FRC) with excellent fracture resistance and non-linear stress-strain curve at ambient temperature. The mechanisms of deformation and fracture in these materials were analyzed in detail, and parallel investigations provided micromechanical models which related the macroscopic behavior to the microstructural parameters [8,9]. However, most of this work was carried out in the ambient temperature regime, where FRC are unlikely to be used. The amount of work on their elevated temperature performance is still limited, and this is especially true for the fracture toughness and fracture resistance. Only a handful of investigations on these topics are available in the open literature [6,10-161. It was demonstrated at the beginning of the 70s It should also be noted that the fracture behavior that high performance C fibers could be successfully of FRC at elevated temperature can be considered incorporated into ceramic glasses to achieve high a litmus test because the presence of a pre-existing strength, tough composite materials [ 141. These notch or crack facilitates the entry of oxygen into early developments were not carried further due to the composite and accelerates any degradation prothe oxidative instability of C fibers at high tempera- cess which might occur. In addition, it is necessary ture. As the new oxidation-resistant SIC fibers de- to find whether the present materials are able to rived by pyrolysis of organometallic precursors achieve the minimum toughness values of became available at the end of this decade, a new 15 MPa.,/E at 1350°C laid down by the industry to generation appeared of glass and glass-ceramic introduce these new materials as structural commatrices reinforced with SIC fibers [5]. While these ponents in gas turbines [17]. The present knowledge 2441
LORCA et al. SiC FIBER-REINFORCED CERAMICS FRC is outlined in this paper. Firstly, the fracture 25 MPa.m from room temperature to 1400cer of the elevated temperature fracture behavior of posite samples exhibited Kic values of toughness and the fracture resistance data at high More recently, Xu et al. [ 15] tested Si3N2-matrix temperature available in the literature are reviewed. SiCt composites in N2 atmosphere. To prepare the These results are completed with new experimental samples, Si3N4 powders were mixed with sintering data obtained in our laboratory on three different additives (Y2O3, Al2O3 and MgO). Layers of pow FRC, and special emphasis is laid on examining the ders and SiC SCS-6 monofilaments were stacked causes of the high temperature degradation and alternatively and hot pressed in a N2 atmosphere their influence on the fracture micromechanisms. Fracture tests were performed by fiexure of notched Finally, the current strategies to improve the high beams, with the notch plane perpendicular to the temperature fracture behavior of FRC are briefly reinforcing fibers. The unidirectional composite discussed exhibited high Kic values from ambient temperature to 1200C; about 36 MPa. m and 50 MPa/m,re- 2. SUMMARY OF PREVIOUS RESULTS spectively for samples with 14% and 29% fraction of fibers when measured as vEJe.Direct Few experimental results have been published on measurement of Kic gave much lower results, 16 the high temperature toughness and crack growth and 10 MPa vm respectively. Xu et al. [15]sus- resistance of FRC, although other mechanical prop- pected that the direct measurements underestimated erties such as load-displacement curves, flexural the composite toughness. Nevertheless, the point in strength or creep behavior have been reported since question is not the absolute values but the lack of the eighties. Such fracture tests are difficult to per- degradation inside the testing temperature interval form and even more difficult to interpret; perform- regardless of the technique used to determine the ing stable tests at high temperatures with toughness precracked samples and extracting from them The results of fracture tests performed in inert meaningful fracture parameters when anisotropy, environments are summarized in Fig. 1. It is nonlinear behavior and inelastic load-displacement remarkable that no detectable degradation of frac- culties, fracture toughness figures and R-curves peratures was loulle n the interval of testing tem- are present is not an easy task. Despite these diffi- ture toughness withi based on linear elastic fracture analysis have beer Fracture tests at high temperatures in air are published. In this respect, the data reported in lit- more recent, and contrary to tests in inert environ- erature must be considered with care. In fact they ments, they showed a slow fracture toughness are usually presented more as a simple way to com- degradation with increasing temperature, Nair and pare homologous materials or to optimize proces- Wang [12] tested a bidirectional woven SiC/Sic sing conditions than as Intrinsic material composite in air. The composite was processed by chemical vapor infiltration of Sic into a woven Farly reported fracture toughness tests at high Nicalon SiC fiber preform. The fiber content was temperatures with SiC-fiber/ ceramic-matrix compo 40% in volume. Fracture tests were performed on sites were performed in inert environments, and compact tension specimens with a notch depth to showed almost constant values from room tempera- sample length of 0.5. KIc values ranged from about ture up to more than 1000.C, Brennan and Prewo [6] measured the fracture properties of sic. fiber glass ceramic composites in Ar atmosphere. TESTS IN INERT ENVIRONMENTS The SiC fiber( Nippon Carbon Co )diameters ran- sI2N,siC [15] The glass matrix was essentially the Corning Glass commercial 9608 lithium-alumino-silicate (LAS). E except that the TiOz nucleating agent laced SI, N, /SIC [15] by a Z O2 nucleating agent. The volume fraction of 2 32887888222445 point bending The 0 /90 cross-plied samples exhi 1 bit Kic values of over 10 MPa,m from ambient 1000°C ernhart et al. [10]and LaImicy et ul. [ll] tested SiC/SiCr composites in an Ar/H2 atmosphere. The LAS Glass/SIC [6 directional flat laminates were manufactured by acking the Nicalon SiC fiber fabrics and then 200 1600 introducing the SiC matrix by a chemical vapor infiltration proccss. Fracture tests also per- Fig. 1. Apparent toughness at initiation as a function formed by flexure of notched beams and the com- temperature in various FRC tested in inert environt
2442 LLORCA et al.: Sic FIBER-REINFORCED CERAMICS of the elevated temperature fracture behavior of FRC is outlined in this paper. Firstly, the fracture toughness and the fracture resistance data at high temperature available in the literature are reviewed. These results are completed with new experimental data obtained in our laboratory on three different FRC, and special emphasis is laid on examining the causes of the high temperature degradation and their influence on the fracture microm~hanisms. Finally, the current strategies to improve the high temperature fracture behavior of FRC are briefly discussed. 2. SUMMARY OF PREVIOUS RESULTS Few experimental results have been published on the high temperature toughness and crack growth resistance of FRC, although other mechanical properties such as load-displacement curves, flexural strength or creep behavior have been reported since the eighties. Such fracture tests are difficult to perform and even more difficult to interpret; performing stable tests at high temperatures with precracked samples and extracting from them meaningful fracture parameters when anisotropy, nonlinear behavior and inelastic load-displacement are present is not an easy task. Despite these difficulties, fracture toughness figures and R-curves based on linear elastic fracture analysis have been published. In this respect, the data reported in literature must be considered with care. In fact they are usually presented more as a simple way to compare homologous materials or to optimize processing conditions than as intrinsic material characteristics. Early reported fracture toughness tests at high temperatures with Sic-fiber/ceramic-matrix composites were performed in inert environments, and showed almost constant values from room temperature up to more than 1000°C. Brennan and Prewo [6] measured the fracture properties of SiCfiber/glass ceramic composites in Ar atmosphere. The SIC fiber (Nippon Carbon Co.) diameters ranged from 5 to 60 pm with an average size of 12 pm. The glass matrix was essentially the Corning Glass commercial 9608 lithium-alumina-silicate (LAS), except that the TiOz nucleating agent was reptaced by a ZrOz nucleating agent. The volume fraction of Sic fibers was about 50%. The fracture tests were carried out on notched beams subjected to threepoint bending. The 0”/90” cross-plied samples exhibit Kit values of over 10 MPa.&ii from ambient temperature to 1000°C. Bernhart ef al. [lo] and Lamicq et al. [l 11 tested SiC/SiCr composites in an Ar/Hz atmosphere. The bidirectional flat laminates were manufactured by stacking the Nicalon SIC fiber fabrics and then introducing the Sic matrix by a chemical vapor infiltration process. Fracture tests were also performed by flexure of notched beams and the composite samples exhibited KIc values of over 25 MPa.6 from room temperature to 1400°C. More recently, Xu et nl. [IS] tested Signs-rnat~x~ SiCr composites in Nz atmosphere. To prepare the samples, S&N4 powders were mixed with sintering additives (YzOs, Al203 and MgO). Layers of powders and SIC SCS-6 monofilaments were stacked alternatively and hot pressed in a N2 atmosphere. Fracture tests were performed by flexure of notched beams, with the notch plane perpendicular to the reinforcing fibers. The unidirectional composite exhibited high KIc values from ambient temperature to 1200°C; about 36 MPa.&ii and 50 MPa.,/iii, respectively for samples with 14% and 29% volume fraction of fibers when measured as m. Direct measurement of KIc gave much lower results, 16 and 10 MPa.& respectively. Xu et al. [15] suspected that the direct measurements underestimated the composite toughness. Nevertheless, the point in question is not the absolute values but the lack of degradation inside the testing temperature interval regardless of the technique used to determine the toughness. The results of fracture tests performed in inert environments are summarized in Fig. 1. It is remarkable that no detectable degradation of fracture toughness within the interval of testing temperatures was found. Fracture tests at high temperatures in air are more recent, and contrary to tests in inert environments, they showed a slow fracture toughness degradation with increasing temperature. Nair and Wang [12] tested a bidirectional woven SiC/SiCr composite in air. The composite was processed by chemical vapor infiltration of Sic into a woven Nicalon SIC fiber preform. The fiber content was 40% in volume. Fracture tests were performed on compact tension specimens with a notch depth to sample length of 0.5. Z&c values ranged from about 0 400 600 1200 1600 Temperature (‘C) Fig. 1. Apparent toughness at initiation as a function of temperature in various FRC tested in inert environments
LLORCA et al.: SiC FIBER-REINFORCED CERAMICS TESTS IN AIR ufactured by stacking several of nicalon (0. 1 um)of pyrolytic C was deposited on the fiber surface, and the sic matrix was introduced into the SIC/SIC [13] preform by the chemical vapor infiltration form of prismatic bars of 10 mm x 3 mm cross-sec tIon Porosity was around 7-8% The second material was an AlyO3 matrix bidirec tionally reinforced with 37 vol. Nicalon SiC fibers. The preform was manufactured by stacking several layers of Nicalon harness satin weave fabric SIC/SIC [12] The fibers were coated sition with a thin layer of BN (20.3 um) and after wards with a thicker layer of Sic (in the range 3- Temperature('C into the preform by a direct metal oxidation process [19]. The composite was received in the Fig. 2. Apparent toughness at initiation as a function of form of prismatic bars of 10 mm x 3 mm cross-sec temperature in various FRC tested in air. 15 MPa m at ambient temperature to 11 MPa: m The third composite was formed by a zrSio4 at 1200.C. Another bidirectional woven composite matrix uniaxially reinforced with 25 vol %SCS-6 produced by chemical vapor infiltration of SiC into SiC monofilaments made by chemical vapor depo- the Nicalon Sic fiber preform was tested by sition of Sic on a C core. Afterwards they were Gomina and Chermant[13]. The volume fraction of coated with a carbon layer of around 3 um in thick- Sic fibers was about 38%. Fracture tests were car- ness, attaining a final diameter of 142 um. The com- ried out by three-point flexure of notched beams. posite was manufactured by aligning SiC filaments Kic values ranged from about 28 MPa m at ambi. and then incorporating the matrix powder around Fareed et al. [14 measured the toughness of out by hot-pressing the matrix and the fibers in a bidirectional woven Al2O3/SiCr composites in air. flowing nitrogen atmosphere [20]. The composite he material was manufactured by direct metal oxi- plates were fully dense and prismatic bars of dation of molten al which was introduced into the 4 mm x. mm cross-section were machined from fiber preform. The volume fraction of Nicalon Sic the plates fibers was about 35%. Fracture tests were per determined by three-point bend tests on notched beams. Kic values ranged from 28 MPa- m at specimens. Notches were machined in the bars with ambient temperature to 15 MPa. m at 1400.C. a thin diamond wire, leading to a notch radius of Figure 2 summarizes all these results of fract 150 Am. The notch length was around 2 mm for the tests in air. It is seen that a gradual degradation first two composites and around 1. 2 mm for the more noticeable from 700C pears as tempera- third one. The notched hars were tested in a cer- ture increases. It is evident. however. that the amic three-point bend testing fixture with either amount of experimental data currently available is 50 mm(for SiC/SiC and Al203/SiC)or 40mm(for y limited, and as a result, there is not conclusive the zrsiO4/SiC) loading span and the specimen evidence of the mechanisms responsible for the placed edgewise on the fixture degradation of the fracture properties of FRC in air The specimen and the fixture, placed in the hi at elevated temperature. This is the subject of the temperature furnace, were loaded through two following sections, where the role of the alumina rods connected to the actuator and to the and fiber degradation in lowering fracture ad cell, respectively, of a servo-mechanical testing ness with temperature is studied in three machine. The external end of one rod was water- FRC cooled to avoid overheating the load cell. The heat- ing rate was 12 C per min and the specimen was 3. MATERIALS AND EXPERIMENTAI test temperature for at least 30 min TECHNIQUES prior to testing. All the tests were performed in air under stroke control, with a cross-head speed of This investigation of the fracture behavior at el- 50 um per min The load (P)and the cross- head dis- evated temperature was performed in three different placement of the testing machine relative to the composites. The first material was made up of a frame( O)were recorded continuously during the SiC matrix bidirectionally (0-90) reinforced with Lests, the latter through a linear-variable differential 35 voL. Nicalon SiC fibers. The preform was man- transducer placed outside the furnace
LLORCA et al.: SiC FIBER-REINFORCED CERAMICS 2443 TESTS IN AIR 40/ 0”” ” ” “‘1 0 A00 600 1200 1600 Temperature (‘C) Fig. 2. Apparent toughness at initiation as a function of temperature in various FRC tested in air. 15 MPa.Jiii at ambient temperature to 11 MPa.6 at 1200°C. Another bidirectional woven composite produced by chemical vapor infiltration of SIC into the Nicalon SIC fiber preform was tested by Gomina and Chermant [13]. The volume fraction of SIC fibers was about 38%. Fracture tests were carried out by three-point flexure of notched beams. Km values ranged from about 28 MPa.,/iii at ambient temperature to 10 MPa.,/iii at 1000°C. Fareed et al. [14] measured the toughness of bidirectional woven AlzOs/SiCr composites in air. The material was manufactured by direct metal oxidation of molten Al which was introduced into the fiber preform. The volume fraction of Nicalon SIC fibers was about 35%. Fracture tests were performed using the chevron notch technique on beams. Krc values ranged from 28 MPa.,/iii at ambient temperature to 15 MPa.6 at 1400°C. Figure 2 summarizes all these results of fracture tests in air. It is seen that a gradual degradation - more noticeable from 700°C - appears as temperature increases. It is evident, however, that the amount of experimental data currently available is very limited, and as a result, there is not conclusive evidence of the mechanisms responsible for the degradation of the fracture properties of FRC in air at elevated temperature. This is the subject of the following sections, where the role of the interface and fiber degradation in lowering fracture toughness with temperature is studied in three different FRC. 3. MATERIALS AND EXPERIMENTAL TECHNIQUES This investigation of the fracture behavior at elevated temperature was performed in three different composites. The first material was made up of a SIC matrix bidirectionally (O”-90”) reinforced with 35 vol.% Nicalon Sic fibers. The preform was manufactured by stacking several layers of Nicalon plain satin weave fabric. A very thin layer (~0.1 pm) of pyrolytic C was deposited on the fiber surface, and the SIC matrix was introduced into the preform by the chemical vapor infiltration process [18]. The composite was received in the form of prismatic bars of 10 mm x 3 mm cross-section. Porosity was around 7-8%. The second material was an A1203 matrix bidirectionally reinforced with 37 vol.% Nicalon SIC fibers. The preform was manufactured by stacking several layers of Nicalon harness satin weave fabric. The fibers were coated by chemical vapor deposition with a thin layer of BN (20.3 pm) and afterwards with a thicker layer of SiC (in the range 3- 4 pm) onto the BN. The matrix was then introduced into the preform by a direct metal oxidation process [19]. The composite was received in the form of prismatic bars of 10 mm x 3 mm cross-section. Porosity was around 7-8%. The third composite was formed by a ZrSiOAmatrix uniaxially reinforced with 2.5 vol.% SCS-6 Sic monofilaments made by chemical vapor deposition of SIC on a C core. Afterwards they were coated with a carbon layer of around 3 pm in thickness, attaining a final diameter of 142 pm. The composite was manufactured by aligning Sic filaments and then incorporating the matrix powder around the uniaxial preform. The consolidation was carried out by hot-pressing the matrix and the fibers in a flowing nitrogen atmosphere [20]. The composite plates were fully dense and prismatic bars of 4 mm x 1.3 mm cross-section were machined from the plates. The fracture behavior of the composites was determined by three-point bend tests on notched specimens. Notches were machined in the bars with a thin diamond wire, leading to a notch radius of 150 pm. The notch length was around 2 mm for the first two composites and around 1.2 mm for the third one. The notched bars were tested in a ceramic three-point bend testing fixture with either 50 mm (for Sic/Sic and A120s/SiC) or 40 mm (for the ZrSiO,/SiC) loading span and the specimen placed edgewise on the fixture. The specimen and the fixture, placed in the high temperature furnace, were loaded through two alumina rods connected to the actuator and to the load cell, respectively, of a servo-mechanical testing machine. The external end of one rod was watercooled to avoid overheating the load cell. The heating rate was 12°C per min and the specimen was held at the test temperature for at least 30 min prior to testing. All the tests were performed in air under stroke control, with a cross-head speed of 50 pm per min. The load (P) and the cross-head displacement of the testing machine relative to the frame (6) were recorded continuously during the tests, the latter through a linear-variable differential transducer placed outside the furnace
LLORCA et al. SiC FIBER-REINFORCED CERAMICS The f avior was obtained from the P-8 have fracture energies around 20 J/ m2, so 5 J/m2 curve using an equivalent elastic crack approach. can be considered an upper limit for the debond This method assumes that the energy release rate of energy. This condition is achieved by coating the the composite during quasi-static crack propagation fibers with thin layers of either C, bn or Mo is equal to that of a linear elastic material which which avoid chemical reactions between the matrix presents the same P-8 curve. The increase in frac- and the fibers during processing, and provide the ture resistance with crack length, R, for the linear weak interfaces with very low fracture energies elastic material can be computed as Experimental observations on different FRC have p2 ac shown that crack deflection and fiber/matrix sliding R (1) give rise to the redistribution of stresses(and thus he dissipation of energy) around a notch or other where B is the specimen thickness and C the speci- strain concentration site by two fundamental mech- men compliance, which is obtained directly from anisms: distributed matrix cracking and fiber fa the P-a curve in an elastic material because there involving pull-out 9]. In the first case, multiple are no residual displacements at zero load. It should mode I matrix cracks grow from the notch, and be noted that this method underestimates the frac- may even extend across the net section prior te ure resistance because the non-linear mechanisms fiber failure If the second mechanism is dominant of energy dissipation are not included. Various the specimen fracture occurs by the propagation of recent studies have shown, however, that the energy a dominant mode I crack from the notch, with fiber release rate and the toughness values provided by failures occurring as the crack extends. Stress redis he equivalent elastic crack approach are compar- tribution is provided by the tractions exerted on the able to those obtained with more sophisticated crack by the failed fibers as they are pulled out methods [21-23]. The accuracy of this approach from the matrix. It should be noted that a third was considered to be sufficient for the purpose of mechanism of energy dissipation by shear damage this investigation, which was to establish the re- in the matrix was found in several C matrix ationship between the microstructural changes composites [9]. However, this damage mode is not occurring during high temperature testing and the analyzed here because these materials are not suit- fracture response of the composites able for high temperature application in oxidizing Once broken, the composites were examined atmospheres sing scanning electron microscopy and energy-dis- Both mechanisms of energy dissipation lead to pensive X-ray microanalysis to determine the micro- the development of a fibrous fracture surface, where structural changes during high temperature the fibers protrude from the matrix. This mor- exposure and the associated failure mechanisms. phology was observed in the fracture surfaces of the The fracture surfaces were first analyzed and the SI C/SiC composite tested at room temperature, as pecimens were then sliced in the longitudinal direc- depicted in Fig. 3(a). a closer examination tion (perpendicular to the fracture surface)with a (Fig 3(b)showed the thin layer of pyrolytic Con low speed diamond saw. The surfaces were polished the lateral fiber surface, indicating that fiber/ matrix successively on diamond cloths of 40, 9, 3 and 1 um debonding took place between the C coating and grain size and finally on alumina of 0. 3 um grain the SiC matrix. The panorama changed completely size. They were cleaned for 30 min by ultrasound in when the material was tested at 1200"C, as is show acetone to remove the alumina from polishing, and in Fig. 3(a). The fracture surfaces at 1200 C wer sputtered with Au-Pd for three minutes before predominantly flat, and the fibers were fractured in being observed in the scanning electron microscope, the crack plane as the matrix crack propagated into them(Fig. 3(C), indicating that fiber/ matrix decohe sion did not take place. This behavior is typical of 4. MICROSTRUCTURAL CHIANGES AND C and Mo coatings which are not stable in oxidiz. ACTURE RESISTANCE ing atmospheres above 700C, and form volatile oxides. The elimination of the coating creates a gap 4.. Interface degi at the interface, which is eventually filled by the for The critical factor to obtain a tough FRC lies in mation of a glassy phase by either matrix or fiber the nature of the fiber matrix interface. If the fibers oxidation, which bonds the fiber to the matrix are strongly bonded to the matrix, a crack Under such conditions, crack deflection cannot take nucleated in the matrix breaks the fibers as it pro- place and the composite fails in a brittle galEs,and the composite is as brittle as the fashion [13, 16 matrix On the contrary, crack deflection and fiber/ The differences in the fracture behavior between matrix sliding at the interface occur when the fibers ambient and elevated temperature are shown in ire weakly bonded to the matrix and this debond Fig. 4, where the fracture resistance is plotted as energy is lower than approximately one fourth of function of the crack length increment, Aa(normal- the fibcr fracture cncrgy [24]. Otherwise the crack izcd by the specimen width, w). The apparent propagates through the fiber, Most ceramic fibers toughness at initiation at 20C was three times
2444 LLORCA et al.: Sic FIBER-REINFORCED CERAMICS The fracture behavior was obtained from the P-6 curve using an equivalent elastic crack approach. This method assumes that the energy release rate of the composite during quasi-static crack propagation is equal to that of a linear elastic material which presents the same P-6 curve. The increase in fracture resistance with crack length, R, for the linear elastic material can be computed as RX: (1) where B is the specimen thickness and C the specimen compliance, which is obtained directly from the P-6 curve in an elastic material because there are no residual displacements at zero load. It should be noted that this method underestimates the fracture resistance because the non-linear mechanisms of energy dissipation are not included. Various recent studies have shown, however, that the energy release rate and the toughness values provided by the equivalent elastic crack approach are comparable to those obtained with more sophisticated methods [21-231. The accuracy of this approach was considered to be sufficient for the purpose of this investigation, which was to establish the relationship between the microstructural changes occurring during high temperature testing and the fracture response of the composites. Once broken, the composites were examined using scanning electron microscopy and energy-dispersive X-ray microanalysis to determine the microstructural changes during high temperature exposure and the associated failure mechanisms. The fracture surfaces were first analyzed and the specimens were then sliced in the longitudinal direction (perpendicular to the fracture surface) with a low speed diamond saw. The surfaces were polished successively on diamond cloths of 40, 9, 3 and 1 pm grain size and finally on alumina of 0.3 pm grain size. They were cleaned for 30 min by ultrasound in acetone to remove the alumina from polishing, and sputtered with Au-Pd for three minutes before being observed in the scanning electron microscope. 4. MICROSTRUCTURAL CHANGES AND FRACTURE RESISTANCE 4.1. Interface degradation The critical factor to obtain a tough FRC lies in the nature of the fiber/matrix interface. If the fibers are strongly bonded to the matrix, a crack nucleated in the matrix breaks the fibers as it propagates, and the composite is as brittle as the matrix. On the contrary, crack deflection and fiber/ matrix sliding at the interface occur when the fibers are weakly bonded to the matrix and this debond energy is lower than approximately one fourth of the fiber fracture energy [24]. Otherwise the crack propagates through the fiber. Most ceramic fibers have fracture energies around 20 J/m*, so 5 J/m2 can be considered an upper limit for the debond energy. This condition is achieved by coating the fibers with thin layers of either C, BN or MO, which avoid chemical reactions between the matrix and the fibers during processing, and provide the weak interfaces with very low fracture energies. Experimental observations on different FRC have shown that crack deflection and fiber/matrix sliding give rise to the redistribution of stresses (and thus the dissipation of energy) around a notch or other strain concentration site by two fundamental mechanisms: distributed matrix cracking and fiber failure involving pull-out [9]. In the first case, multiple mode I matrix cracks grow from the notch, and may even extend across the net section prior to fiber failure. If the second mechanism is dominant, the specimen fracture occurs by the propagation of a dominant mode I crack from the notch, with fiber failures occurring as the crack extends. Stress redistribution is provided by the tractions exerted on the crack by the failed fibers as they are pulled out from the matrix. It should be noted that a third mechanism of energy dissipation by shear damage in the matrix was found in several C matrix composites [9]. However, this damage mode is not analyzed here because these materials are not suitable for high temperature application in oxidizing atmospheres. Both mechanisms of energy dissipation lead to the development of a fibrous fracture surface, where the fibers protrude from the matrix. This morphology was observed in the fracture surfaces of the Sic/Sic composite tested at room temperature, as depicted in Fig. 3(a). A closer examination (Fig. 3(b)) showed the thin layer of pyrolytic C on the lateral fiber surface, indicating that fiber/matrix debonding took place between the C coating and the Sic matrix. The panorama changed completely when the material was tested at 12OO”C, as is shown in Fig. 3(a). The fracture surfaces at 1200°C were predominantly flat, and the fibers were fractured in the crack plane as the matrix crack propagated into them (Fig. 3(c)), indicating that fiber/matrix decohesion did not take place. This behavior is typical of C and MO coatings, which are not stable in oxidizing atmospheres above =7OO”C, and form volatile oxides. The elimination of the coating creates a gap at the interface, which is eventually filled by the formation of a glassy phase by either matrix or fiber oxidation, which bonds the fiber to the matrix. Under such conditions, crack deflection cannot take place and the composite fails in a brittle fashion [13,16]. The differences in the fracture behavior between ambient and elevated temperature are shown in Fig. 4, where the fracture resistance is plotted as function of the crack length increment, Aa (normalized by the specimen width, I+‘). The apparent toughness at initiation at 20°C was three times
LlORCA et al. siC FIBER-REINFORCED CERAMICS 2445 Fig 3. (a)Fracture surfaces of the SiC/SiCr composite at 20C (left)and 1200 c (right).(b) SiC fiber pulled out from the matrix. The pyrolytic C coating on the fiber surface is marked with an arrow.(c) SiC fibers bonded to the matrix and broken in the crack plane during fracture at 1200.c higher than at 1200C, and in addition, the interface produced an almost fat R-curve, the frac resistance increased rapidly with crack ture resistance similar to that of the matrix ambient temperature, as more energy is di dependent of the crack in the crack wake through fiber pull-out The degradation mechanism presented above contrary,the high temperature embrittlement be partially suppressed by using coatings with better induced by the development of a strong fiber/ matrix oxidation resistance, such as BN. The oxidation of
higher ambie in the contra induce LLORCA et al.: Sic FIBER-REINFORCED CERAMICS 2445 Fig. 3. (a) Fracture surfaces of the SiC/SiCr composite at 20°C (left) and 1200°C (right). (b) Sic fiber pulled out from the matrix. The pyrolytic C coating on the fiber surface is marked with an arrow. (c) Sic fibers bonded to the matrix and broken in the crack plane during fracture at 1200°C. than at 12OO”C, and in addition, the fracture interface produced an almost flat R-curve, the fracnce increased rapidly with crack length at ture resistance being similar to that of the matrix nt temperature, as more energy is dissipated and almost independent of the crack length. crack wake through fiber pull-out. On the The degradation mechanism presented above can .ry, the high temperature embrittlement be partially suppressed by using coatings with better :d by the development of a strong fiber/matrix oxidation resistance, such as BN. The oxidation of