composes Part A: applied science and manufacturing ELSEVIER Composites: Part A 33(2002)1209-1218 www.elsevier.com/locate/composites fracture behaviour of cross-ply Nicalon/CAS-II glass-ceramic matrix composite laminate at room and elevated temperatures A.Yasmin"P. bowen igh Performance Applications, University of Birmingham, Edgbaston, Birmingham B15 2TT, UK Received 23 January 2002; revised 23 May 2002: accepted 26 June 2002 Abstract The fracture behaviour of cross-ply Nicalon/CAS-nI glass ceramic matrix co mposite laminate has been investigated as a function of temperature, loading rate and environment. Tests were carried out at 20, 600, 800, 1000C in air and also at 1000.C in vacuum. Increased flexural strength was obtained with increased loading rate but it decreased with increasing test temperature. Although the change in flexural rength was not very significant for the loading rates(10-and 10-mm/min)employed in this study except at 600.Cin air, the influence of environment was dramatic. The flexural strength at 1000C in vacuum was comparable to the value obtained at room temperature in air, however, the flexural strength at 1000 C in air reduced significantly from its room temperature value. Two major damage modes have been identified: fibrous at room temperature and non-fibrous at high temperature. c 2002 Elsevier Science Ltd. All rights reserved Keywords: A Ceramic-matrix composites(CMCs); B. Fracture; D. Mechanical testing: D. Electron microscopy 1. Introduction advance [4-7]. For example, by employing continuous fibres the fracture toughness of CMCs or G-CMCs can be When combined with high specific strength and stiffness, increased to =30 MPavm [8] compared to monolithic the ability of ceramic materials to withstand severe ceramics that show only about 3 MPam[4]. In addition environments including heat, abrasion and high oxidation they retain load carrying capacity after matrix cracks are compared to conventional materials have made these initiated whereas monolithic ceramics show catastrophic materials particularly suitable for the development of high failure temperature structural components such as aero-engine, gas After achieving excellent properties from unidirectional turbine and nuclear reaction furnaces [1, 2]. Consequently, composites, the next interest goes to cross-ply or angle-ply they have the potential to produce damage tolerant and laminates due to their potential to be applied under tough materials that can be used at temperatures as high as multiaxial stress conditions. An extensive research on the 1000C for several 100 h. However. the main drawback of damage mechanism of cross-ply G-CMCs under uniaxial these materials lies in their reduced ductility and fracture nsile loading condition is available in the literature toughness, which makes ceramic components potentially [9-11]. It shows that the presence of off-axis plies changes prone to catastrophic failure 3]. Therefore, in the last two the failure mechanisms. There are two modes of decades, considerable attention has been given to the cracking in the cross-ply laminates: one is transverse development of continuous silicon carbide fibre reinforced cracking in the 90 plies and the other is matrix cracking in ceramic( CMCs)and glass-ceramic matrix composites he 0 plies. Therefore, the failure processes start at lower stresses than that for the unidirectional composite (G-CMCs)to provide excellent toughness by providing However, the presence of transverse cracks in the cross- various energy dissipating processes (i.e. matrix micro- cracking, fibre/matrix debonding, fibre pullout) via load ply laminate has only a limited effect on the evolution of matrix cracks and hence, the composite strength is transfer and crack deflection mechanisms during crack controlled by the o ply proportional limit point where orresponding autho address: Centre for Intelligent matrix cracking starts, and not by the 90 ply failure point Processing of Composite #330, Northwestern University However, the room temperature behaviour of cross-ply Evanston. IL 60208 USA. 491-7961;fax:+1-847-491-5227 laminates may deteriorate at high temperatur E-mail address: a-yasmi western.edu(A. Yasmin) the early cracking of 90 plies, which may penetrate past 1359-835X/02/S-see front matter e 2002 Elsevier Science Ltd. All rights reserved. PI:S1359-835X(02)000799
Fracture behaviour of cross-ply Nicalon/CAS-II glass–ceramic matrix composite laminate at room and elevated temperatures A. Yasmin*, P. Bowen School of Metallurgy and Materials/IRC in Materials for High Performance Applications, University of Birmingham, Edgbaston, Birmingham B15 2TT, UK Received 23 January 2002; revised 23 May 2002; accepted 26 June 2002 Abstract The fracture behaviour of cross-ply Nicalon/CAS-II glass–ceramic matrix composite laminate has been investigated as a function of temperature, loading rate and environment. Tests were carried out at 20, 600, 800, 1000 8C in air and also at 1000 8C in vacuum. Increased flexural strength was obtained with increased loading rate but it decreased with increasing test temperature. Although the change in flexural strength was not very significant for the loading rates (1023 and 1021 mm/min) employed in this study except at 600 8C in air, the influence of environment was dramatic. The flexural strength at 1000 8C in vacuum was comparable to the value obtained at room temperature in air, however, the flexural strength at 1000 8C in air reduced significantly from its room temperature value. Two major damage modes have been identified: fibrous at room temperature and non-fibrous at high temperature. q 2002 Elsevier Science Ltd. All rights reserved. Keywords: A. Ceramic–matrix composites (CMCs); B. Fracture; D. Mechanical testing; D. Electron microscopy 1. Introduction When combined with high specific strength and stiffness, the ability of ceramic materials to withstand severe environments including heat, abrasion and high oxidation compared to conventional materials have made these materials particularly suitable for the development of high temperature structural components such as aero-engine, gas turbine and nuclear reaction furnaces [1,2]. Consequently, they have the potential to produce damage tolerant and tough materials that can be used at temperatures as high as 1000 8C for several 100 h. However, the main drawback of these materials lies in their reduced ductility and fracture toughness, which makes ceramic components potentially prone to catastrophic failure [3]. Therefore, in the last two decades, considerable attention has been given to the development of continuous silicon carbide fibre reinforced ceramic (CMCs) and glass–ceramic matrix composites (G–CMCs) to provide excellent toughness by providing various energy dissipating processes (i.e. matrix microcracking, fibre/matrix debonding, fibre pullout) via load transfer and crack deflection mechanisms during crack advance [4–7]. For example, by employing continuous fibres the fracture toughness of CMCs or G–CMCs can be increased to $30 MPapm [8] compared to monolithic ceramics that show only about 3 MPapm [4]. In addition, they retain load carrying capacity after matrix cracks are initiated whereas monolithic ceramics show catastrophic failure. After achieving excellent properties from unidirectional composites, the next interest goes to cross-ply or angle-ply laminates due to their potential to be applied under multiaxial stress conditions. An extensive research on the damage mechanism of cross-ply G–CMCs under uniaxial tensile loading condition is available in the literature [9–11]. It shows that the presence of off-axis plies changes the failure mechanisms. There are two major modes of cracking in the cross-ply laminates: one is transverse cracking in the 908 plies and the other is matrix cracking in the 08 plies. Therefore, the failure processes start at lower stresses than that for the unidirectional composite. However, the presence of transverse cracks in the crossply laminate has only a limited effect on the evolution of matrix cracks and hence, the composite strength is controlled by the 08 ply proportional limit point where matrix cracking starts, and not by the 908 ply failure point. However, the room temperature behaviour of cross-ply laminates may deteriorate at high temperature due to the early cracking of 908 plies, which may penetrate past 1359-835X/02/$ - see front matter q 2002 Elsevier Science Ltd. All rights reserved. PII: S1 35 9 -8 35 X( 02 )0 0 07 9 -9 Composites: Part A 33 (2002) 1209–1218 www.elsevier.com/locate/compositesa * Corresponding author. Present address: Centre for Intelligent Processing of Composites, Room #330, Northwestern University, Evanston, IL 60208, USA. Tel.: þ1-847-491-7961; fax: þ1-847-491-5227. E-mail address: a-yasmin@northwestern.edu (A. Yasmin)
A. Yasmin, P. Bowen/Composites: Part A 33(2002)1209-1218 the fibres of the 0 plies and expose them to environmental An optical micrograph of the cross-section of cross-ply attack Nicalon/CAS is shown in Fig. 1. Nicalon/CAS consists of The high strength and high toughness values obtained at aIcium-alumino-silicate (CA amic matrIx room temperature are believed to be due to the presence of a reinforced with 35% continuous silicon carbide fibre weakly bonded carbon-rich interfacial layer developed(Nicalon) by volume. Specimens with dimensions of about during hot pressing [12-14]. However, at higher tempera- 50 mm long, 5 mm wide and 3. 7 mm thick were cut from the ture the carbon interface layer is removed by oxidation composite plate at O to the plate axis using diamond tooling reaction upon matrix cracking and a very strong silica bond fracture tests were performed on as-received and plane is developed [15-17]. This silica bond increases interfacial sided specimens under three-point bending with a total shear strength [13, 18-20] and causes failure of the fibres loading span of 40 mm. The specimens were edge-loaded along with the matrix instead of any type of crack deflecting and the schematic of this loading configuration is given in mechanism. The degradation of strength of fibres by thermal Fig. 2. The three-point bend technique allows simple decomposition may be another reason of embrittlement; specimen geometry to be used under a relatively simple however, it is reported by others [21] that it is less crucial for fixture and also utilization of small quantity of plate material Nicalon/CAS glass-ceramic matrix composite system. The very efficiently. Tests were carried out at room temperature embrittlement process of cross-ply Nicalon/CAS (9 is 600, 800, 1000C in air and also at 1000C in vacuum found to be similar to unidirectional Nicalon/CAS except Crosshead displacement rates of 10 and 10 mm/min 90 plies crack prior to 0 ply matrix cracking; however, were used in this study. For all high temperature tests, embrittlement does not occur until the matrix in the 0 plies specimens were heated at a rate of 10 C/min and held at the cracks. Therefore, if these are the materials to be used at required temperature for I h under a low load of 10n to high temperature, a detailed investigation on their behaviour allow the system to reach equilibrium. An Instron 8501 under loading is important servohydraulic machine equipped with a 5 kN load cell was The present work illustrates the effects of temperature sed for fracture tests in air. whereas vacuum tests were and loading rate on the failure mechanism of cross-ply carried out in a further Instron 8501 servohydraulic machine Nicalon/CAS glass-ceramic matrix composite laminate equipped with a vacuum chamber and a load cell of 10 kN under three-point bending subjected to monotonic loading During the test, a load versus crosshead displacement to failure. the environmental effect on the fracture record was taken for all tests and the nominal maximum behaviour has also been investigated. In addition, stress or flexural strength was calculated using the following comparison has been drawn between the fracture behaviour formula of unidirectional and cross-ply Nicalon/CAS glass-ceramic 1) cre P is the maximum load, L is the half of the loading 2. Experimental B is the specimen thickness and w is the specimen width. This is called nominal maximum stress. since the The material used in this study is a cross-ply [0/90J4s stresses in the O and 90 plies are not the same and also shift icalon/CAS-Il glass-ceramic matrix composite laminates in proportion as damage takes place manufactured by Corning Industries, USA. The composite The fracture surfaces of the specimens were investigated was supplied as a plate of dimensions 150 X 150 x 3.7 mm. using a Hitachi $4000 Field Emission Gun(FEG)scanning electron microscope(SEM) with an accelerating voltage of 4 kv, which did not cause any charging of the specimen and 0°fibr 4 250 Fig. 1. Optical micrograph of the n of as-received cross-ply s all dimensions are in mm Nicalon/CAS-ll glass-ceramic matrix composite. Fig. 2. A schematic of three-point bend loading configurati
the fibres of the 08 plies and expose them to environmental attack. The high strength and high toughness values obtained at room temperature are believed to be due to the presence of a weakly bonded carbon-rich interfacial layer developed during hot pressing [12–14]. However, at higher temperature the carbon interface layer is removed by oxidation reaction upon matrix cracking and a very strong silica bond is developed [15–17]. This silica bond increases interfacial shear strength [13,18–20] and causes failure of the fibres along with the matrix instead of any type of crack deflecting mechanism. The degradation of strength of fibres by thermal decomposition may be another reason of embrittlement; however, it is reported by others [21] that it is less crucial for Nicalon/CAS glass–ceramic matrix composite system. The embrittlement process of cross-ply Nicalon/CAS [9] is found to be similar to unidirectional Nicalon/CAS except 908 plies crack prior to 08 ply matrix cracking; however, embrittlement does not occur until the matrix in the 08 plies cracks. Therefore, if these are the materials to be used at high temperature, a detailed investigation on their behaviour under loading is important. The present work illustrates the effects of temperature and loading rate on the failure mechanism of cross-ply Nicalon/CAS glass–ceramic matrix composite laminate under three-point bending subjected to monotonic loading to failure. The environmental effect on the fracture behaviour has also been investigated. In addition, a comparison has been drawn between the fracture behaviour of unidirectional and cross-ply Nicalon/CAS glass–ceramic matrix composite laminates. 2. Experimental The material used in this study is a cross-ply [0/90]4s Nicalon/CAS-II glass–ceramic matrix composite laminates manufactured by Corning Industries, USA. The composite was supplied as a plate of dimensions 150 £ 150 £ 3.7 mm3 . An optical micrograph of the cross-section of cross-ply Nicalon/CAS is shown in Fig. 1. Nicalon/CAS consists of calcium–alumino-silicate (CAS) glass–ceramic matrix reinforced with 35% continuous silicon carbide fibre (Nicalon) by volume. Specimens with dimensions of about 50 mm long, 5 mm wide and 3.7 mm thick were cut from the composite plate at 08 to the plate axis using diamond tooling. Fracture tests were performed on as-received and plane sided specimens under three-point bending with a total loading span of 40 mm. The specimens were edge-loaded and the schematic of this loading configuration is given in Fig. 2. The three-point bend technique allows simple specimen geometry to be used under a relatively simple fixture and also utilization of small quantity of plate material very efficiently. Tests were carried out at room temperature, 600, 800, 1000 8C in air and also at 1000 8C in vacuum. Crosshead displacement rates of 1023 and 1021 mm/min were used in this study. For all high temperature tests, specimens were heated at a rate of 10 8C/min and held at the required temperature for 1 h under a low load of 10 N to allow the system to reach equilibrium. An Instron 8501 servohydraulic machine equipped with a 5 kN load cell was used for fracture tests in air, whereas vacuum tests were carried out in a further Instron 8501 servohydraulic machine equipped with a vacuum chamber and a load cell of 10 kN. During the test, a load versus crosshead displacement record was taken for all tests and the nominal maximum stress or flexural strength was calculated using the following formula s ¼ 3PL BW2 ð1Þ where P is the maximum load, L is the half of the loading span, B is the specimen thickness and W is the specimen width. This is called nominal maximum stress, since the stresses in the 0 and 908 plies are not the same and also shift in proportion as damage takes place. The fracture surfaces of the specimens were investigated using a Hitachi S4000 Field Emission Gun (FEG) scanning electron microscope (SEM) with an accelerating voltage of 4 kV, which did not cause any charging of the specimen and Fig. 1. Optical micrograph of the cross-section of as-received cross-ply Nicalon/CAS-II glass–ceramic matrix composite. Fig. 2. A schematic of three-point bend loading configuration. 1210 A. Yasmin, P. Bowen / Composites: Part A 33 (2002) 1209–1218
A. Yasmin, P. Bowen /Composites: Part A 33(2002)1209-1218 thereby permitted scanning without gold coating 图 Ramp rate,0001 mm/min Occasionally, the pull-out length of fibres was measured B Ramp rate, 0.1 mm/min from SEM FEG micrographs taken at a tilt angle of 4. Fig 3 shows the variation of flexural strength of cross-ply 8"900 Nicalon/CAs glass-ceramic matrix composite with tem- perature, (in air)at the two different loading rates of 10-3 300 and 10 mm/min. From this figure, it is apparent that there 2 200 increased from room temperature to 1000C in air, and 800 is a steady decrease of flexural strength as the temperature is approximately 41, 32 and 26% of the room temperature value is obtained at 600, 800 and 1000C, respectively. lus, the temperature effect is highly significant and this gnificant decrease in flexural strength with temperature is 20°C,air1000°c,air1000°c,vac. true for both loading rates (10 and 10mm/min) Fig. 4. Effects of environment and temperature on the flexural streng employed in this study. However, the influence of loading cross-ply Nicalon/CAS glass-ceramic matrix composite rate for a given test temperature was not found to be pronounced except at 600C where a large degree of variation of flexural strength was obtained at the higher The load versus displacement curves for all test loading rate temperatures at both loading rates (10-and In Fig. 4, the flexural strength of cross-ply Nicalon/CAS 10 mm/min) are summarised in Fig. 5(a)and(b), at 1000C in vacuum has been compared with the values respectively. It is found that as the temperature increases obtained at both room temperature and 1000C in air. The oth peak load and the corresponding displacement flexural strength values at 1000C in vacuum are found to decreases and two types of failure modes can be identified be 7-14% higher than the room temperature values for the from these curves loading rates employed in this study. However, the flexural trength values at 1000 C in air are found to be significantl 1. Non-catastrophic or fibrous: tests at 20C in air and lower, just 22-26% of the room temperature and 20-23% 1000C in vacuum show non-linear behaviour until the of the 1000c in vacuum values. this observation maximum load is reached, which is typical of damage therefore, suggests a strong influence of environment on tolerant materials. After reaching the peak value the load the flexural strength of Nicalon/CAS at high temperatures drops gradually but in a discontinuous manner. Finally the specimens fail by gross collapse Catastrophic or non-fibrous: tests at 600, 800 and 1000C in air show linear behaviour until the maximum Ramp rate, 0.001 mm/mim load is reached and then fail catastrophically Since the area under the load-deflection curve is proportional to the ability of the material to absorb energy ng testing, Fig. 5(a)and(b) also energy involved in non-catastrophic failure specimens is 5300 higher than the catastrophic failure specimens ig. 6(a) illustrates an SEM micrograph of the side face of a specimen, loaded monotonically to failure at a ading rate of 10 mm/min at room temperature in air There are no dominant mode -I cracks and cracking is diffuse. Damage consists of multiple matrix microcrack ing, debonding and fibre pullout. The corresponding load- displacement curve in Fig. 5(a) shows non-linearity to a 0 200 400 600 800 1000 1200 displacement value of 0.6 mm to reach the peak load (780N). The load then drops to a residual value(240 N, The variation of flexural strength of cross-ply Nicalon/CAS glass 0.67 mm), which is about 30%o of the peak load and the ic matrix composite in air with temperature at two different loadin long tail afterwards is attributed to extensive pullout of he individual fibres. The load-displacement curve
thereby permitted scanning without gold coating. Occasionally, the pull-out length of fibres was measured from SEM FEG micrographs taken at a tilt angle of 458. 3. Results Fig. 3 shows the variation of flexural strength of cross-ply Nicalon/CAS glass–ceramic matrix composite with temperature (in air) at the two different loading rates of 1023 and 1021 mm/min. From this figure, it is apparent that there is a steady decrease of flexural strength as the temperature is increased from room temperature to 1000 8C in air, and approximately 41, 32 and 26% of the room temperature value is obtained at 600, 800 and 1000 8C, respectively. Thus, the temperature effect is highly significant and this significant decrease in flexural strength with temperature is true for both loading rates (1023 and 1021 mm/min) employed in this study. However, the influence of loading rate for a given test temperature was not found to be pronounced except at 600 8C where a large degree of variation of flexural strength was obtained at the higher loading rate. In Fig. 4, the flexural strength of cross-ply Nicalon/CAS at 1000 8C in vacuum has been compared with the values obtained at both room temperature and 1000 8C in air. The flexural strength values at 1000 8C in vacuum are found to be 7–14% higher than the room temperature values for the loading rates employed in this study. However, the flexural strength values at 1000 8C in air are found to be significantly lower, just 22–26% of the room temperature and 20–23% of the 1000 8C in vacuum values. This observation, therefore, suggests a strong influence of environment on the flexural strength of Nicalon/CAS at high temperatures. The load versus displacement curves for all test temperatures at both loading rates (1023 and 1021 mm/min) are summarised in Fig. 5(a) and (b), respectively. It is found that as the temperature increases both peak load and the corresponding displacement decreases and two types of failure modes can be identified from these curves: 1. Non-catastrophic or fibrous: tests at 20 8C in air and 1000 8C in vacuum show non-linear behaviour until the maximum load is reached, which is typical of damage tolerant materials. After reaching the peak value the load drops gradually but in a discontinuous manner. Finally, the specimens fail by gross collapse. 2. Catastrophic or non-fibrous: tests at 600, 800 and 1000 8C in air show linear behaviour until the maximum load is reached and then fail catastrophically. Since the area under the load–deflection curve is proportional to the ability of the material to absorb energy during testing, Fig. 5(a) and (b) also clearly show that the energy involved in non-catastrophic failure specimens is higher than the catastrophic failure specimens. Fig. 6(a) illustrates an SEM micrograph of the side face of a specimen, loaded monotonically to failure at a loading rate of 1023 mm/min at room temperature in air. There are no dominant mode-I cracks and cracking is diffuse. Damage consists of multiple matrix microcracking, debonding and fibre pullout. The corresponding load– displacement curve in Fig. 5(a) shows non-linearity to a displacement value of 0.6 mm to reach the peak load (780 N). The load then drops to a residual value (240 N, 0.67 mm), which is about 30% of the peak load and the long tail afterwards is attributed to extensive pullout of the individual fibres. The load–displacement curve, Fig. 4. Effects of environment and temperature on the flexural strength of cross-ply Nicalon/CAS glass–ceramic matrix composite. Fig. 3. The variation of flexural strength of cross-ply Nicalon/CAS glass– ceramic matrix composite in air with temperature at two different loading rates. A. Yasmin, P. Bowen / Composites: Part A 33 (2002) 1209–1218 1211
A. Yasmin, P. Bowen/ Composites: Part A 33(2002)1209-1218 Tensile side 1000 000°C,vac 600°C,air 750pm 100C, air Tensile side 0.00.20.40.60.81.0121.41.6 1000 0°,vac 400三 750pu Non-Fibrous 600°C,ai ig. 6. SEM micrographs of the side face of specimens subjected to 800°C,air monotonic failure at room temperature Loading rate (a)10 mm/min and 00c (b)10 mm/min 10 mm/min(Fig 3). Fig. 7(b) shows the fracture surface 0.00.20.4060.81.01.21.41.6 of the mid-section of the specimen that showed the lowest Displacement, (mm) peak load(332.5 N) among all specimens tested at the higher loading rate(10 mm/min). The fractograph shows Fig. 5. Load-displacement curves for all temperatures employed in this significant pullout compared to the specimen tested at the udy Loading rate(a)10-mm/min and(b)10-mm/min lower loading rate, and the average pullout length is approximately 50 um. The corresponding load-displace therefore, indicates a gross collapse and this type of ment curve in Fig. 5(b) also shows that the peak load occurs failure mechanism is usually referred to as ' fibrous'in at a higher displacement (0.27 mm) compared to the nature specimen tested at the lower loading rate, 10 mm/min However,a dominant crack and fibre bridging are After that the load drops stepwise with a higher displace found as the loading rate is increased to 10 mm/min as ment, perhaps an indication of fibre failure as bundles shown in Fig. 6(b). The corresponding load-displacement Fig. 7(e) illustrates the side face of the specimen that curve in Fig. 5(b) also shows similar behaviour to th showed the highest peak load(636.2 N) among all speci load-displacement curve at the lower loading rate mens tested at the higher loading rate of 10 mm/min 10-mm/min), however, exhibits a residual load which (Note that, the load-displacement curve for this test is not is about 30% of the peak load at a higher displacement available due to a problem with the chart recorder). In this (275N, 0.9 mm) than that observed at the lower loading case, the specimen shows a tortuous mode-I crack with random fibre failure and the fibre pullout length is about ig. 7(a) shows the fracture surface of the mid-section of 135 um. Therefore, this type of fracture behaviour can be a specimen tested at 600C in air with a loading rate of characterised as'mixed and a transition from non-fibrous to 10 mm/min. The 0 fibre layers of the fracture surface are mixed fracture occurs as the loading rate is increased at found nearly flat with limited fibre pullout. The correspond- 600C in air g load-displacement curve in Fig. 5(a) also shows Fig &(a)and(b) show SEM micrographs of the side faces catastrophic behaviour. The specimen loses its load carrying of the specimens tested at a loading rate of 10 mm/min at capacity(95%)instantly after the peak load(293N) is 800 and 1000C in air, respectively. Damage consists of reached. However, the specimens show more scatter in mode-I cracking in both conditions. However, specimens flexural strength values as the loading rate is increased to tested at 800 C in air show occasional presence of a little o
therefore, indicates a gross collapse and this type of failure mechanism is usually referred to as ‘fibrous’ in nature. However, a dominant crack and fibre bridging are found as the loading rate is increased to 1021 mm/min as shown in Fig. 6(b). The corresponding load–displacement curve in Fig. 5(b) also shows similar behaviour to the load–displacement curve at the lower loading rate (1023 mm/min), however, exhibits a residual load which is about 30% of the peak load at a higher displacement (275 N, 0.9 mm) than that observed at the lower loading rate, 1023 mm/min. Fig. 7(a) shows the fracture surface of the mid-section of a specimen tested at 600 8C in air with a loading rate of 1023 mm/min. The 08 fibre layers of the fracture surface are found nearly flat with limited fibre pullout. The corresponding load–displacement curve in Fig. 5(a) also shows catastrophic behaviour. The specimen loses its load carrying capacity (,95%) instantly after the peak load (293 N) is reached. However, the specimens show more scatter in flexural strength values as the loading rate is increased to 1021 mm/min (Fig. 3). Fig. 7(b) shows the fracture surface of the mid-section of the specimen that showed the lowest peak load (332.5 N) among all specimens tested at the higher loading rate (1021 mm/min). The fractograph shows significant pullout compared to the specimen tested at the lower loading rate, and the average pullout length is approximately 50 mm. The corresponding load–displacement curve in Fig. 5(b) also shows that the peak load occurs at a higher displacement (0.27 mm) compared to the specimen tested at the lower loading rate, 1023 mm/min. After that the load drops stepwise with a higher displacement, perhaps an indication of fibre failure as bundles. Fig. 7(c) illustrates the side face of the specimen that showed the highest peak load (636.2 N) among all specimens tested at the higher loading rate of 1021 mm/min. (Note that, the load–displacement curve for this test is not available due to a problem with the chart recorder). In this case, the specimen shows a tortuous mode-I crack with random fibre failure and the fibre pullout length is about 135 mm. Therefore, this type of fracture behaviour can be characterised as ‘mixed’ and a transition from non-fibrous to mixed fracture occurs as the loading rate is increased at 600 8C in air. Fig. 8(a) and (b) show SEM micrographs of the side faces of the specimens tested at a loading rate of 1021 mm/min at 800 and 1000 8C in air, respectively. Damage consists of mode-I cracking in both conditions. However, specimens tested at 800 8C in air show occasional presence of a little 08 Fig. 6. SEM micrographs of the side face of specimens subjected to monotonic failure at room temperature. Loading rate (a) 1023 mm/min and (b) 1021 mm/min. Fig. 5. Load–displacement curves for all temperatures employed in this study. Loading rate (a) 1023 mm/min and (b) 1021 mm/min. 1212 A. Yasmin, P. Bowen / Composites: Part A 33 (2002) 1209–1218
A. Yasmin, P. Bowen /Composites: Part A 33(2002)1209-1218 1213 Tensile side 90 ply 0°ply 231um Tensile side 90°py 273pm Tensile side Fig. 8. SEM micrographs of the side face of specimens subjected to monotonic failure at a loading rate of 10 mm/min in air(a)800C and(b) 1000°C. load and therefore, the specimen can maintain a high residual strength for a higher displacement value (218 N, 1.1mm) Table 1 summarises the flexural strength values and failure modes for all test conditions applied in this study pecimens tested at room temperature and 1000C in vacuum show fibrous failure, whereas specimens tested at Fig. 7. SEM f the specimens subjected to monotonic failure at ate:10-3mm/min;(b)loadin 800 and 1000C in air show non-fibrous failure at both 600°Cin 0.1 mm/min showed the lowest flexural strength; (c) loading loading rates employed in this study. A transition from non- min; specimen showed the highest flexural strength. fibrous to mixed failure is observed at 600C when the loading rate is increased from 10 to 10 mm/min. the above observation, therefore, signify that the flexural fibre pullout but specimens tested at 1000C in air show strength of cross-ply Nicalon/CAS is strongly dependent absolutely no fibre pullout. In general, in both cases the on temperature and environment, and the loading rate can fracture surfaces are flat and the corresponding loa also be important. displacement curves in Fig. 5(b) show catastrophic failure at a very low peak load and displacement values. However,a higher flexural strength is obtained at 800C than at 1000C. Tensile side This type of failure is usually referred to as 'non-fibrous However a fibrous failure can be obtained at 1000C uum as shown in Fig. 9. The failure is characterised by multiple transverse matrix microcracking, fibre debonding, fibre bridging and fibre pullout. This type of failure mechanism is similar to that observed in roor temperature tests and consequently a high flexural strengt is obtained which is even higher than the room temperature values. The correspo ponding load-displacement curve Fig. 9. SEM micrograph of the side face of a specimen subjected to Fig 5(a) shows that the load drops gradually after the peak monotonic failure at 1000C in vacuum at a loading rate of 10-mm/min
fibre pullout but specimens tested at 1000 8C in air show absolutely no fibre pullout. In general, in both cases the fracture surfaces are flat and the corresponding load– displacement curves in Fig. 5(b) show catastrophic failure at a very low peak load and displacement values. However, a higher flexural strength is obtained at 800 8C than at 1000 8C. This type of failure is usually referred to as ‘non-fibrous’ failure. However, a fibrous failure can be obtained at 1000 8C, but in vacuum as shown in Fig. 9. The failure is characterised by multiple transverse matrix microcracking, fibre debonding, fibre bridging and fibre pullout. This type of failure mechanism is similar to that observed in room temperature tests and consequently a high flexural strength is obtained which is even higher than the room temperature values. The corresponding load–displacement curve in Fig. 5(a) shows that the load drops gradually after the peak load and therefore, the specimen can maintain a high residual strength for a higher displacement value (218 N, 1.1 mm). Table 1 summarises the flexural strength values and failure modes for all test conditions applied in this study. Specimens tested at room temperature and 1000 8C in vacuum show fibrous failure, whereas specimens tested at 800 and 1000 8C in air show non-fibrous failure at both loading rates employed in this study. A transition from non- fibrous to mixed failure is observed at 600 8C when the loading rate is increased from 1023 to 1021 mm/min. The above observation, therefore, signify that the flexural strength of cross-ply Nicalon/CAS is strongly dependent on temperature and environment, and the loading rate can also be important. Fig. 8. SEM micrographs of the side face of specimens subjected to monotonic failure at a loading rate of 1021 mm/min in air (a) 800 8C and (b) 1000 8C. Fig. 7. SEM micrographs of the specimens subjected to monotonic failure at 600 8C in air. (a) Loading rate: 1023 mm/min; (b) loading rate: 0.1 mm/min, specimen showed the lowest flexural strength; (c) loading rate, 1021 mm/min; specimen showed the highest flexural strength. Fig. 9. SEM micrograph of the side face of a specimen subjected to monotonic failure at 1000 8C in vacuum at a loading rate of 1023 mm/min. A. Yasmin, P. Bowen / Composites: Part A 33 (2002) 1209–1218 1213
Yasmin, P. Bowen/ Composites: Part A 33(2002)1209-1218 through the fibres [24] and inhibits the energy dissipating of fracture test results performed on cross-ply Nicalon/CAS mechanisms such as debonding crack deflection along the Flexural strength values Failure modes 0 plies, fibre pullout and crack bridging responsible for the (average)(MPa) toughening behaviour in Nicalon/CAS. Therefore, the lesser amount and/or shorter length of fibre pullout are an Loading rate Loading rate Loading rate loading ra indication of stronger interfacial bonding strength. For mm/min mm/m mm/min) mm/min) shorter fibre pullout than the tests at 20C in air(Fig. 6)and 20(air)489,507(498)497,531(514) Fibrous Fibrous eventually an intermediate flexural strength is obtained 600air)191,2190205)215,316,427Non- fibrous Mixed However, the tests at 800 and 1000C in air(Fig 8)exhibit almost no fibre pullout and finally a very low flexural 800 air) 162, 143 (162) 180, 195 (184) Non-ibrous Non-ibrous strength is obtained. Furthermore, with increasing tempera- 1000(vac.)561,577(569)549,553(551) Fibrous Fibrous ir, the density of ma cracking is foun decrease and a total absence of multiple cracks is observed at 1000C in air( Fig 8(b). The load-displacement curves related to non-fibrous behaviour in Fig. 5 also exhibit 4. Discussion linearity and low displacement to reach the maximum load and drop instantly from the maximum load to almost zero load similar to that observed in monolithic ceramics [25] 4.1. Efects of temperature The single dominant crack and low flexural strength at high temperature have also been characterised by Allen [20] by Fibrous failure, as observed in room temperature and bending tests on unidirectional Nicalon/CAS and by Puente 1000C in vacuum, consists of debonding at the fibre- et al. [26] by tensile tests on cross-ply Nicalon/CAS matrix interface, multiple matrix cracking, fibre pullout and fibre bridging. This type of failure mechanism is 4.2. Efects of loading rate attributed to the presence of a well-defined carbon layer at he fibre-matrix interface developed during fabrication In the present study, a strong dependence of flexural [12-14]. This carbon layer is strong enough to transfer the strength on the loading rate is observed only at 600C in load from the matrix to the fibre, however, weak enough to A transition from non -fibrous to mixed failure occurs when il preferentially prior to fibre failure and ahead of an the loading rate increases from 103 to 10-mm/min advancing matrix crack [22]. The carbon layer, therefore, Consistently, a higher flexural strength (about 57%)is ncreases flaw tolerance by allowing the debonding to obtained at the higher loading rate. However, the scatter of occur prior to crack propagation through the fibres and/or flexural strength values only at 600C at the higher ramp by deflecting matrix cracks parallel to the fibres [23]. The ate is not very clear, but it should be mentioned that this is clean or smooth fibre pullout surface, observed in this the temperature where transition in failure mechanism study at room temperature(Fig. 6(b) and at 1000C in occurs and can be controlled by simply changing the loadin vacuum(Fig. 9) is an indication of sufficiently low rate. It is also observed that at this higher loading rate as the interfacial bond strength between the fibre and the matrix. pullout length increases the flexural strength for the The corresponding load-displacement curves in Fig. 5(a) corresponding specimen increases. The strong influence of and(b)also shows non-linearity and high displacement loading rate on the fracture behaviour of unidirectional values to reach the maximum load and afterwards a long Nicalon/CAS has also been reported by Allen et al. [27] tail with a high residual load capability. All these They observed a fibrous failure at 600C in air when the characteristics together signify the damage tolerance and loading rate was increased to 10 mm/min but mixed and high toughness nature of the composite. Hence, a high fibrous failures even at 800C when the loading rates were flexural strength is obtained both at room temperature in air increased to 1 and 10 mm/min, respectively. It can and at1000° in vacuun. therefore, be suggested that a fibrous failure might result On the contrary, the high temperature tests in air show a even at high temperature if a sufficiently high loading rate is decrease in flexural strength. Oxidation of interfacial carbon employed. layer can start at temperature as low as 600c by Furthermore, the reasons for the higher flexural strength atmospheric oxygen ingress into the bulk of the material and mixed failure at the higher loading rate, 10 mm/min upon matrix cracking and becomes virulent at temperatures compared to the lower loading rate at 600C in air can be above 800C[12]. The oxidation process is followed by the explained as less time was available for 'silica bridge formation of ' silica bridge reaction at the fibre formation in the former case. At the lower loading rate, the interface, which produces a very high interfacial bond test took more than 4 h to fail; therefore, there was sufficient strength [13, 18-20]. When the bond strength exceeds a time for silica bridging reaction, whereas at the higher certain value, matrix cracks in the 0 plies can penetrate loading rate the specimen took only 5 min to fail. Therefore
4. Discussion 4.1. Effects of temperature Fibrous failure, as observed in room temperature and 1000 8C in vacuum, consists of debonding at the fibre– matrix interface, multiple matrix cracking, fibre pullout and fibre bridging. This type of failure mechanism is attributed to the presence of a well-defined carbon layer at the fibre–matrix interface developed during fabrication [12–14]. This carbon layer is strong enough to transfer the load from the matrix to the fibre, however, weak enough to fail preferentially prior to fibre failure and ahead of an advancing matrix crack [22]. The carbon layer, therefore, increases flaw tolerance by allowing the debonding to occur prior to crack propagation through the fibres and/or by deflecting matrix cracks parallel to the fibres [23]. The clean or smooth fibre pullout surface, observed in this study at room temperature (Fig. 6(b)) and at 1000 8C in vacuum (Fig. 9) is an indication of sufficiently low interfacial bond strength between the fibre and the matrix. The corresponding load–displacement curves in Fig. 5(a) and (b) also shows non-linearity and high displacement values to reach the maximum load and afterwards a long tail with a high residual load capability. All these characteristics together signify the damage tolerance and high toughness nature of the composite. Hence, a high flexural strength is obtained both at room temperature in air and at 1000 8C in vacuum. On the contrary, the high temperature tests in air show a decrease in flexural strength. Oxidation of interfacial carbon layer can start at temperature as low as 600 8C by atmospheric oxygen ingress into the bulk of the material upon matrix cracking and becomes virulent at temperatures above 800 8C [12]. The oxidation process is followed by the formation of ‘silica bridge’ reaction at the fibre–matrix interface, which produces a very high interfacial bond strength [13,18–20]. When the bond strength exceeds a certain value, matrix cracks in the 08 plies can penetrate through the fibres [24] and inhibits the energy dissipating mechanisms such as debonding, crack deflection along the 08 plies, fibre pullout and crack bridging responsible for the toughening behaviour in Nicalon/CAS. Therefore, the lesser amount and/or shorter length of fibre pullout are an indication of stronger interfacial bonding strength. For example, the tests at 600 8C in air (Fig. 7) show fewer and shorter fibre pullout than the tests at 20 8C in air (Fig. 6) and eventually an intermediate flexural strength is obtained. However, the tests at 800 and 1000 8C in air (Fig. 8) exhibit almost no fibre pullout and finally a very low flexural strength is obtained. Furthermore, with increasing temperature in air, the density of matrix cracking is found to decrease and a total absence of multiple cracks is observed at 1000 8C in air (Fig. 8(b)). The load–displacement curves related to non-fibrous behaviour in Fig. 5 also exhibit linearity and low displacement to reach the maximum load and drop instantly from the maximum load to almost zero load similar to that observed in monolithic ceramics [25]. The single dominant crack and low flexural strength at high temperature have also been characterised by Allen [20] by bending tests on unidirectional Nicalon/CAS and by Puente et al. [26] by tensile tests on cross-ply Nicalon/CAS. 4.2. Effects of loading rate In the present study, a strong dependence of flexural strength on the loading rate is observed only at 600 8C in air. A transition from non-fibrous to mixed failure occurs when the loading rate increases from 1023 to 1021 mm/min. Consistently, a higher flexural strength (about 57%) is obtained at the higher loading rate. However, the scatter of flexural strength values only at 600 8C at the higher ramp rate is not very clear, but it should be mentioned that this is the temperature where transition in failure mechanism occurs and can be controlled by simply changing the loading rate. It is also observed that at this higher loading rate as the pullout length increases the flexural strength for the corresponding specimen increases. The strong influence of loading rate on the fracture behaviour of unidirectional Nicalon/CAS has also been reported by Allen et al. [27]. They observed a fibrous failure at 600 8C in air when the loading rate was increased to 1021 mm/min, but mixed and fibrous failures even at 800 8C when the loading rates were increased to 1 and 10 mm/min, respectively. It can, therefore, be suggested that a fibrous failure might result even at high temperature if a sufficiently high loading rate is employed. Furthermore, the reasons for the higher flexural strength and mixed failure at the higher loading rate, 1021 mm/min compared to the lower loading rate at 600 8C in air can be explained as less time was available for ‘silica bridge’ formation in the former case. At the lower loading rate, the test took more than 4 h to fail; therefore, there was sufficient time for silica bridging reaction, whereas at the higher loading rate the specimen took only 5 min to fail. Therefore, Table 1 Summary of fracture test results performed on cross-ply Nicalon/CAS Temp. (8C) Flexural strength values (average) (MPa) Failure modes Loading rate (1023 mm/min) Loading rate (1021 mm/min) Loading rate (1021 mm/min) Loading rate (1021 mm/min) 20 (air) 489, 507 (498) 497, 531 (514) Fibrous Fibrous 600 (air) 191, 219 (205) 215, 316, 427 (319) Non-fibrous Mixed 800 (air) 162, 162 (162) 180, 195 (187) Non-fibrous Non-fibrous 1000 (air) 107, 155 (131) 111, 121 (115) Non-fibrous Non-fibrous 1000 (vac.) 561, 577 (569) 549, 553 (551) Fibrous Fibrous 1214 A. Yasmin, P. Bowen / Composites: Part A 33 (2002) 1209–1218
A. Yasmin, P. Bowen/Composites: Part A 33(2002)1209-1218 1215 the rate of silica bridging reaction at the fibre-matrix fibres in a lithium -alumino-silicate matrix but only when interface is important. the specimen was stressed beyond the matrix microcrackin stress level 4.3. Effects of environment At high temperature in vacuum, there is very little 4.4. Effects of fibre orientation oxygen present and therefore, no oxidation of carbon layer and no'silica bridge reaction at the interface. The intact A comparison in flexural strength values between carbon layer at the interface allows a progressive fracture unidirectional [20] and cross-ply Nicalon/CAS containin 35 volume fraction of fibres under identical conditions has mechanism similar to the room temperature tests in the bsence of environmental attack at the fibre-matrix been drawn in Table 2. from this table. it is found that at a interface. As a result, a high flexural strength is obtained loading rate of 10 mm/min the cross-ply composite shows at high temperature in vacuum [28, 29). However, the 7 and about 70 and 52% of the flexural strength of unidirectional 49 higher flexural strength of cross-ply Nicalon/CAS at composite at room temperature(also at 1000C in vacuum) 1000C in vacuum compared to the tests carried out at and high temperature, respectively. However, at a loadin room temperature in air at 10 and 10 mm/min loading rate of 10 mm/min, the cross-ply shows about 63 and 42%o rates,respectively, indicate the possibility of silica of the flexural strength of unidirectional composite at room bridging reaction even in vacuum. Since the temperature temperature (also at 1000C in vacuum) and high gh and the test takes long time to he excess temperature, respectively. Therefore, the effect is more oxygen still available in the matrix or fibre itself may pronounced at high temperature rather than at room diffuse through the matrix and heal any faw on the fibre temperature. Furthermore, at 600C in air and under a surface by silica coating or transfer a greater proportion of constant loading rate range(10 to 10 mm/min)the stress to intact matrix material via silica bridging [30]. The cross-ply Nicalon/CAS shows the non-fibrous to mixed er flexural strength at the lower loading ra failure transition. whereas the unidirectional Nicalon/CAS compared to the higher loading rate can be attributed to the shows the non-fibrous to fibrous failure transition [27]. All increased duration of time in the former case ns could be attributed to the different Therefore, the diffusion of oxygen through the matrix for a composite architecture, the effective volume fraction of longer period simply increases the possibility of silica reinforcing fibres and the damage mechanisms in these bridging reaction. The improved flexural strength composites. Nicalon/CAs due to silica bridging reaction at consider The lower flexural strength of cross-ply composite ably higher temperature(800C)is also reported by ompared to the unidirectional composite at room ssett et al. 301 temperature could be attributed to the lower effective From the above observations, it is important to volume fraction of fibres in the former case. If it is assumed mention that high temperature alone (in the that the 90 plies carry no load and also have no influence oxygen environment) does not deteriorate the strength of Nicalon fibre and/or the carbon interface and subsequently the strength of the composite. Ishikawa 31] has confirmed that the as-received Nicalon fibres can retain A comparison in flexural strength values between unidirectional and cros excellent mechanical properties at 1000C for 1000 h in ply Nicalon/CAS composites air but starts to degrade at temperatures above 1300C. However, some mechanical degradation of Nicalon fibres Average flexural strength(MPa) Ocross-ply always occurs during composite fabrication as reported elsewhere [28 Nicalon/CAS(oui) Nicalon/CAS(across It is also reported that the environment does not cause any embrittlement at high temperature(1200C)until the Loading rate, 10mm/mi matrix has been cracked [13]. In this regard, the thermal 20(air) exposure of specimens at high temperature(600-1000Cin 600(air) 800(air) air) during the holding period is expected not to cause any 1000(air) serious damage to the specimens. If there 1000(vacuum) bridging reaction, it should be only to the outer regions near Loading rate, 10-mm/min the surface and not to the interior of the sample since there 20(air) 0.66 are no matrix cracks for oxygen penetration into the bulk 600(air) the specimen. This thermal exposure, therefore, has no 800(air) apparent influence on the flexural strength of the composite. 1000(vacuum) 0.4 925 551 Mah et al. [28] have also reported similar behaviour. They found high temperature oxygen embrittlement of Nicalon Data from Ref [20]
the rate of silica bridging reaction at the fibre–matrix interface is important. 4.3. Effects of environment At high temperature in vacuum, there is very little oxygen present and therefore, no oxidation of carbon layer and no ‘silica bridge’ reaction at the interface. The intact carbon layer at the interface allows a progressive fracture mechanism similar to the room temperature tests in the absence of environmental attack at the fibre–matrix interface. As a result, a high flexural strength is obtained at high temperature in vacuum [28,29]. However, the 7 and 14% higher flexural strength of cross-ply Nicalon/CAS at 1000 8C in vacuum compared to the tests carried out at room temperature in air at 1021 and 1023 mm/min loading rates, respectively, indicate the possibility of silica bridging reaction even in vacuum. Since the temperature is high and the test takes long time to fail, the excess oxygen still available in the matrix or fibre itself may diffuse through the matrix and heal any flaw on the fibre surface by silica coating or transfer a greater proportion of stress to intact matrix material via silica bridging [30]. The 14% higher flexural strength at the lower loading rate compared to the higher loading rate can be attributed to the increased duration of testing time in the former case. Therefore, the diffusion of oxygen through the matrix for a longer period simply increases the possibility of silica bridging reaction. The improved flexural strength of Nicalon/CAS due to silica bridging reaction at considerably higher temperature (,800 8C) is also reported by Blissett et al. [30]. From the above observations, it is important to mention that high temperature alone (in the absence of oxygen environment) does not deteriorate the strength of Nicalon fibre and/or the carbon interface and subsequently the strength of the composite. Ishikawa [31] has confirmed that the as-received Nicalon fibres can retain excellent mechanical properties at 1000 8C for 1000 h in air but starts to degrade at temperatures above 1300 8C. However, some mechanical degradation of Nicalon fibres always occurs during composite fabrication as reported elsewhere [28]. It is also reported that the environment does not cause any embrittlement at high temperature (,1200 8C) until the matrix has been cracked [13]. In this regard, the thermal exposure of specimens at high temperature (600–1000 8C in air) during the holding period is expected not to cause any serious damage to the specimens. If there is any silica bridging reaction, it should be only to the outer regions near the surface and not to the interior of the sample since there are no matrix cracks for oxygen penetration into the bulk of the specimen. This thermal exposure, therefore, has no apparent influence on the flexural strength of the composite. Mah et al. [28] have also reported similar behaviour. They found high temperature oxygen embrittlement of Nicalon fibres in a lithium–alumino-silicate matrix but only when the specimen was stressed beyond the matrix microcracking stress level. 4.4. Effects of fibre orientation A comparison in flexural strength values between unidirectional [20] and cross-ply Nicalon/CAS containing 35% volume fraction of fibres under identical conditions has been drawn in Table 2. From this table, it is found that at a loading rate of 1023 mm/min the cross-ply composite shows about 70 and 52% of the flexural strength of unidirectional composite at room temperature (also at 1000 8C in vacuum) and high temperature, respectively. However, at a loading rate of 1021 mm/min, the cross-ply shows about 63 and 42% of the flexural strength of unidirectional composite at room temperature (also at 1000 8C in vacuum) and high temperature, respectively. Therefore, the effect is more pronounced at high temperature rather than at room temperature. Furthermore, at 600 8C in air and under a constant loading rate range (1023 to 1021 mm/min) the cross-ply Nicalon/CAS shows the non-fibrous to mixed failure transition, whereas the unidirectional Nicalon/CAS shows the non-fibrous to fibrous failure transition [27]. All these observations could be attributed to the different composite architecture, the effective volume fraction of reinforcing fibres and the damage mechanisms in these composites. The lower flexural strength of cross-ply composite compared to the unidirectional composite at room temperature could be attributed to the lower effective volume fraction of fibres in the former case. If it is assumed that the 908 plies carry no load and also have no influence Table 2 A comparison in flexural strength values between unidirectional and crossply Nicalon/CAS composites Temp. (8C) Average flexural strength (MPa) scross-ply /suni Unidirectional Nicalon/CASa (suni) Cross-ply Nicalon/CAS (scross-ply) Loading rate,1023 mm/min 20 (air) 693 498 0.71 600 (air) 420 205 0.48 800 (air) 305 162 0.53 1000 (air) 233 131 0.56 1000 (vacuum) 805 569 0.70 Loading rate,1021 mm/min 20 (air) 778 514 0.66 600 (air) 685 322 0.47 800 (air) 450 187 0.41 1000 (air) 288 115 0.40 1000 (vacuum) 925 551 0.60 a Data from Ref. [20]. A. Yasmin, P. Bowen / Composites: Part A 33 (2002) 1209–1218 1215
A. Yasmin, P. Bowen/ Composites: Part A 33(2002)1209-1218 on the strength of the 0 plies, the number of effective plies (or 0 plies) taking part in carrying the load is half of the otal number of plies in the composite. Consequently, the ideal nominal strength of cross-ply will be just 50% of Region A the unidirectional composite. Consistent with this assump- Blissett et al. [ 30] have ted the flexural strength Neutral axis of as-received cross-ply Nicalon/CAs is approximately 50% that of unidirectional composite. However, the Region B 63-70% recovery of flexural strength at room temperature in the case of cross-ply compared to the unidirectional composite indicates some role of 90 plies in this study. An Fig. 10. Schematic of the cross-section of [0/90]4s Nicalon/CAS co analytical model discussed in Section 4.5 also confirms beam with the strain distribution through the thickness The significant drop in flexural strength of laminate (50-60%0) compared to the unidirectional lami nate at high temperature should be attributed to both some damages in the 90 plies of region B, and equilibrium effective volume fraction of fibres and the different failure of forces. it can be written mechanisms In the case of cross-ply, the microcracking in d 90 plies known as'transverse cracking,, always occurs at h lower stresses than matrix cracking in the o plies [9-11 32). It is of interest to note that cyclic loading on cross-ply where h is the height of the specimen, d is the distance of Nicalon/CAS at 800C in air shows fatigue damage below neutral axis from the maximum tensile stress surface, E, is a peak stress of 110 MPa, but rapid failure occurs at a peak the elastic modulus of unidirectional laminate in the stress of 110 MPa [33]. This stress is in close agreement longitudinal direction, E2 is the elastic modulus with the matrix cracking stress or flexural strength unidirectional laminate in the transverse direction. ra is (113 MPa)of cross-ply Nicalon/CAs at 1000C in the stiffness reduction factor due to cracking in the 90 plies, obtained in this study(Fig. 5). This is also close to the reported matrix microcracking stress value obtained by EA is the maximum strain in region A and EB is the maximum strain in region B tensile testing [13]. Therefore, the transverse cracking in the 90 plies should occur at lower stresses and it can be as Assuming that both unidirectional and cross-ply beams ave the same geometry and both fail when the maximum low as 60 MPa as detected by acoustic emission technique flexural stress in region B, OB, reaches the tensile strength 34]. At room temperature the composite strength is controlled by the 0 ply proportional limit where matrix cracking starts and not by the 90 ply failure. Therefore, OB(o)=Fit good toughness behaviour similar to the unidirectional Then the ratio of ultimate loads of cross-ply and composite is obtained. On the contrary, at high tempera- unidirectional composites can be written ture the earlier occurrence of cracks in 90 plies complicate the failure process. In addition to the lower Pcross-ply-h stress, the cracking in the 90 plies also occurs at a d strain. about 0.05%0 than the first matrix cracki unidirectional composite, 0. 1% [34]. This crack penetrate past the fibres of the 0 plies and expose them to environmental attack at elevated temperature. Therefore, (4) there is an earlier removal of the interfacial carbon layer In order to compare flexural strengths, Eqs. (1)and(4)give very low flexural strength in cross-ply compared to the h unidirectional composite 4.5. Analytical modelling of the flexural strength of cross- E3 ply Nicalon/CAS under three-point bending d)+(1+n2E Fig. 10 shows the schematic of the cross-section of [0/90J4s Nicalon/CAS composite beam with the strain dis- Table 3 shows the calculated cross-ply/uni ratios at tribution through the thickness, where regions A and B are different stiffness reduction factor(r2) value(where E on and tension, respectively. Considering 127 GPa, E2= 112 GPa, from Ref [35])and they are in linear beam theory and assuming no damage in region A and close agreement with the measured values given in
on the strength of the 08 plies, the number of effective plies (or 08 plies) taking part in carrying the load is half of the total number of plies in the composite. Consequently, the ‘ideal’ nominal strength of cross-ply will be just 50% of the unidirectional composite. Consistent with this assumption, Blissett et al. [30] have reported the flexural strength of as-received cross-ply Nicalon/CAS is approximately 50% that of unidirectional composite. However, the 63–70% recovery of flexural strength at room temperature in the case of cross-ply compared to the unidirectional composite indicates some role of 908 plies in this study. An analytical model discussed in Section 4.5 also confirms this. The significant drop in flexural strength of cross-ply laminate (50–60%) compared to the unidirectional laminate at high temperature should be attributed to both effective volume fraction of fibres and the different failure mechanisms. In the case of cross-ply, the microcracking in 908 plies known as ‘transverse cracking’, always occurs at lower stresses than matrix cracking in the 08 plies [9–11, 32]. It is of interest to note that cyclic loading on cross-ply Nicalon/CAS at 800 8C in air shows fatigue damage below a peak stress of 110 MPa, but rapid failure occurs at a peak stress of 110 MPa [33]. This stress is in close agreement with the matrix cracking stress or flexural strength (,113 MPa) of cross-ply Nicalon/CAS at 1000 8C in air obtained in this study (Fig. 5). This is also close to the reported matrix microcracking stress value obtained by tensile testing [13]. Therefore, the transverse cracking in the 908 plies should occur at lower stresses and it can be as low as 60 MPa as detected by acoustic emission technique [34]. At room temperature the composite strength is controlled by the 08 ply proportional limit where matrix cracking starts and not by the 908 ply failure. Therefore, good toughness behaviour similar to the unidirectional composite is obtained. On the contrary, at high temperature the earlier occurrence of cracks in 908 plies complicate the failure process. In addition to the lower stress, the cracking in the 908 plies also occurs at a lower strain, about 0.05%, than the first matrix cracking in unidirectional composite, 0.1% [34]. This crack may penetrate past the fibres of the 08 plies and expose them to environmental attack at elevated temperature. Therefore, there is an earlier removal of the interfacial carbon layer and the onset of silica bridge reaction, which produces a very low flexural strength in cross-ply compared to the unidirectional composite. 4.5. Analytical modelling of the flexural strength of crossply Nicalon/CAS under three-point bending Fig. 10 shows the schematic of the cross-section of [0/90]4s Nicalon/CAS composite beam with the strain distribution through the thickness, where regions A and B are under compression and tension, respectively. Considering linear beam theory and assuming no damage in region A and some damages in the 908 plies of region B, and equilibrium of forces, it can be written d h ¼ 1 1 þ ffiffiffiffiffiffiffiffiffi E1þr2E2 E1þE2 q ð2Þ where h is the height of the specimen, d is the distance of neutral axis from the maximum tensile stress surface, E1 is the elastic modulus of unidirectional laminate in the longitudinal direction, E2 is the elastic modulus of unidirectional laminate in the transverse direction, r2 is the stiffness reduction factor due to cracking in the 908 plies, 1A is the maximum strain in region A and 1B is the maximum strain in region B. Assuming that both unidirectional and cross-ply beams have the same geometry and both fail when the maximum flexural stress in region B, sB, reaches the tensile strength, F1t sBð08Þ ¼ F1t ð3Þ Then the ratio of ultimate loads of cross-ply and unidirectional composites can be written as Pcross-ply Puni ¼ h d 1 þ E2 E1 1 2 d h 3 þ 1 þ r2 E2 E1 d h 3 " # ð4Þ In order to compare flexural strengths, Eqs. (1) and (4) give scross-ply suni ¼ h d 1 þ E2 E1 1 2 d h 3 þ 1 þ r2 E2 E1 d h 3 " # ð5Þ Table 3 shows the calculated scross-ply/suni ratios at different stiffness reduction factor (r2) value (where E1 ¼ 127 GPa; E2 ¼ 112 GPa; from Ref. [35]) and they are in close agreement with the measured values given in Fig. 10. Schematic of the cross-section of [0/90]4s Nicalon/CAS composite beam with the strain distribution through the thickness. 1216 A. Yasmin, P. Bowen / Composites: Part A 33 (2002) 1209–1218
A. Yasmin, P. Bowen /Composites: Part A 33(2002)1209-1218 1217 Table 3 Acknowledgements Calculated value of across-plyouni ratios at different r2 values One of the authors(AY) would like to acknowledge the 2-000 financial supports from the Overseas Research Scholarship 0.578 (ORS), UK and the School of Metallurgy and Materials University of Birmingham, UK. Thanks are due to rolls Royce Plc., Derby for the provision of materials and Dr TJA 0.65 Doel for his thoughtful discussions on this work ever needed. Thanks are also due to dr jj luo and dr jl abot for advices in preparing the manuscrip 0.3 0.7 0.73 0.533 References [1] Donald LW, McMillan PW. Ceramic-matrix composites. J Mater Sci Table 2. If r2 is equal to zero, there is no contribution 1976:11:949-72. om the 90 plies and the strength of cross-ply is only [2] Okamura K. Ceramic-matri Mater1995:4(3):247-59 57%that of unidirectional composite. This is probably 3] Johnson DR, Stiegler JO. Structural cer th 50% recovery of flexural strength at high temperature in 4] Ruhle M, Evans AG. High toughness ceramics and ceramic air indicates the possible damages in the 0 plies due composites. Prog Mater Sci 1989: 33: 85-167. [5] Herbell TP. Eckel AG. Ceramic matrix composites for rocket engine earlier cracking in the 90" plies and the onset of silica turbine applications. Trans ASME 1993: 115: 64-9 bridging reaction. However, the measured 60-70% 6] Ohnabe H, Masaki S, Onozuka M. Miyahara K. Sasa T. Potential recovery of flexural strength found at room temperature application of ceramic matrix composites to aero-engine composites. and also at high temperature in vacuum at both loading rates corresponds to r2 values between 0. 1 and 0.35. It is [7] Mason S Minor R. Razzell AG. Ceramic matrix composites for aero- engines. Mater World 1993 Jan: 16-18 ported that a typical value of r2 is 0.25 [36] and this [8] Allen RF, Beevers C, Bowen P. Fracture and fatigue of a Nicalon/ corresponds to a calculated 67%o recovery of the strength CAS continuous fibre-reinforced glass-ceramic matrix composite of unidirectional composite Composites, Part A 1993: 24(2): 150-6 concluded that the value of r2 is not zero at room [9] Kahraman R, Mandell JF, Deibert MC. High temperature mechanical temperature and also at high temperature in vacuum. As behaviour of multidirectional Nicalon/CAS-ll composite. J Mater Sci 1995:30:6329-38 result, the 90 plies contribute to the strength of cross- [10 Lee SS. Stinchcomb ww Damage mechanisms and fracture modes in Nicalon/CAS-Il laminates. Key Engng Mater 1996: 121-122 227-56. [11] Zawada LP, Butkus LM, Hartman GA. Tensile and fatigue behaviour of silicon carbide fibre reinforced aluminosilicate glass. J Am Ceram 5. Conclusions Soc1991;74(11):2851-8 [12] Bleay SM, Scott VD. The relationship between interface microstruc The flexural strength of cross-ply Nicalon/CAS glass- ture and bulk properties in glass-ceramic matrix composites. Euromat’9l1991;2:300-6 eramic matrix composite is found to be strongly dependent [13] Lewis MH, Daniel AM. Chamberlain A, Pharaoh MW, Cain MG on test temperature and environment. The higher the test Microstructure-property relationships in silicate-matrix composites. temperature in air, the lower the flexural strength of the J Microsc1993;1692):109-18. laminate for both loading rates employed in this study. This [14] Cooper RF, Chyung K. Structure and chemistry of fibre-matrix interfaces in SiC fibre- reinforced glas has been deduced to be due to the silica bridge reaction at electron micre tudy. J Mater Sci 1987: 22: 3148-60 high temperature in air, which results in changes in failure [15] Huger M, Fargeot D, Gault C Ultrasonic characterization of oxidation mechanisms from fibrous at room temperature to non- mechanisms in Nicalon/C/SiC composites. J Am Ceram Soc 1994 fibrous at higher temperature. No influence of loading rate is 77(10:2554-60 found both at high and low temperatures but a moderate [16] Filipuzzi L, Camus G, Naslain R Oxidation mechanisms and kinetics influence is found at an intermediate temperature, 600C, of ID-SiC/C/SiC composite materials. 1. An experimental approach. J Am Ceram Soc1994;77(2):459-66 where a transition from non-fibrous to mixed failure occurs [17] Filipuzzi L, Naslain R. Oxidation mechanisms and kinetics of ID. with increased loading rate. However, fracture tests at SiC/C/SiC composite materials. Il Modeling. J Am Ceram Soc 1994 1000C in vacuum show fibrous behaviour and high 77(2)467-80 flexural strength similar to the room temperature tests and [18] Grande DH, Mandell JF, Hong KCC. Fibre-matrix bond strength studies of glass, ceramic and metal matrix composites. J Mater Sci are thus consistent with oxygen promoting silica-bridging 1988;23:311-28 reactions at high temperature in ai [19] Tiwan A. Modelling mechanical response of SiC/CAS-lI ceramic
Table 2. If r2 is equal to zero, there is no contribution from the 908 plies and the strength of cross-ply is only 57% that of unidirectional composite. This is probably the case at high temperature in air. However, the 40– 50% recovery of flexural strength at high temperature in air indicates the possible damages in the 08 plies due to earlier cracking in the 908 plies and the onset of silica bridging reaction. However, the measured 60–70% recovery of flexural strength found at room temperature and also at high temperature in vacuum at both loading rates corresponds to r2 values between 0.1 and 0.35. It is reported that a typical value of r2 is 0.25 [36] and this corresponds to a calculated 67% recovery of the strength of unidirectional composite. Therefore, it can be concluded that the value of r2 is not zero at room temperature and also at high temperature in vacuum. As a result, the 908 plies contribute to the strength of crossply composite as observed in the present study. 5. Conclusions The flexural strength of cross-ply Nicalon/CAS glass– ceramic matrix composite is found to be strongly dependent on test temperature and environment. The higher the test temperature in air, the lower the flexural strength of the laminate for both loading rates employed in this study. This has been deduced to be due to the ‘silica bridge’ reaction at high temperature in air, which results in changes in failure mechanisms from fibrous at room temperature to non- fibrous at higher temperature. No influence of loading rate is found both at high and low temperatures but a moderate influence is found at an intermediate temperature, 600 8C, where a transition from non-fibrous to mixed failure occurs with increased loading rate. However, fracture tests at 1000 8C in vacuum show fibrous behaviour and high flexural strength similar to the room temperature tests and are thus consistent with oxygen promoting silica-bridging reactions at high temperature in air. Acknowledgements One of the authors (AY) would like to acknowledge the financial supports from the Overseas Research Scholarship (ORS), UK and the School of Metallurgy and Materials, University of Birmingham, UK. Thanks are due to Rolls Royce Plc., Derby for the provision of materials and Dr TJA Doel for his thoughtful discussions on this work ever needed. Thanks are also due to Dr JJ Luo and Dr JL Abot for advices in preparing the manuscript. References [1] Donald LW, McMillan PW. Ceramic–matrix composites. J Mater Sci 1976;11:949–72. [2] Okamura K. Ceramic–matrix composites (CMC). Adv Compos Mater 1995;4(3):247–59. [3] Johnson DR, Stiegler JO. Structural ceramics R&D. Adv Mater Process 1990;9:55–61. [4] Ruhle M, Evans AG. High toughness ceramics and ceramic composites. Prog Mater Sci 1989;33:85–167. [5] Herbell TP, Eckel AG. Ceramic matrix composites for rocket engine turbine applications. Trans ASME 1993;115:64–9. [6] Ohnabe H, Masaki S, Onozuka M, Miyahara K, Sasa T. Potential application of ceramic matrix composites to aero-engine composites. Composites, Part A 1999;30:489–96. [7] Mason S, Minor RJ, Razzell AG. Ceramic matrix composites for aeroengines. Mater World 1993;Jan:16–18. [8] Allen RF, Beevers CJ, Bowen P. Fracture and fatigue of a Nicalon/ CAS continuous fibre-reinforced glass–ceramic matrix composite. Composites, Part A 1993;24(2):150–6. [9] Kahraman R, Mandell JF, Deibert MC. High temperature mechanical behaviour of multidirectional Nicalon/CAS-II composite. J Mater Sci 1995;30:6329–38. [10] Lee SS, Stinchcomb WW. Damage mechanisms and fracture modes in Nicalon/CAS-II laminates. Key Engng Mater 1996;121–122: 227–56. [11] Zawada LP, Butkus LM, Hartman GA. Tensile and fatigue behaviour of silicon carbide fibre reinforced aluminosilicate glass. J Am Ceram Soc 1991;74(11):2851–8. [12] Bleay SM, Scott VD. The relationship between interface microstructure and bulk properties in glass–ceramic matrix composites. Euromat’91 1991;2:300–6. [13] Lewis MH, Daniel AM, Chamberlain A, Pharaoh MW, Cain MG. Microstructure–property relationships in silicate–matrix composites. J Microsc 1993;169(2):109–18. [14] Cooper RF, Chyung K. Structure and chemistry of fibre–matrix interfaces in SiC fibre-reinforced glass–ceramic composites: an electron microscopy study. J Mater Sci 1987;22:3148–60. [15] Huger M, Fargeot D, Gault C. Ultrasonic characterization of oxidation mechanisms in Nicalon/C/SiC composites. J Am Ceram Soc 1994; 77(10):2554–60. [16] Filipuzzi L, Camus G, Naslain R. Oxidation mechanisms and kinetics of 1D-SiC/C/SiC composite materials. 1. An experimental approach. J Am Ceram Soc 1994;77(2):459–66. [17] Filipuzzi L, Naslain R. Oxidation mechanisms and kinetics of 1DSiC/C/SiC composite materials. II. Modeling. J Am Ceram Soc 1994; 77(2):467–80. [18] Grande DH, Mandell JF, Hong KCC. Fibre–matrix bond strength studies of glass, ceramic and metal matrix composites. J Mater Sci 1988;23:311–28. [19] Tiwan A. Modelling mechanical response of SiC/CAS-II ceramic Table 3 Calculated value of scross-ply/suni ratios at different r2 values r2 d/h scross-ply/suni 0 0.578 0.57 0.05 0.573 0.59 0.1 0.568 0.61 0.15 0.563 0.63 0.2 0.558 0.65 0.25 0.554 0.67 0.3 0.549 0.69 0.35 0.545 0.71 0.4 0.541 0.73 0.45 0.537 0.75 0.5 0.533 0.76 A. Yasmin, P. Bowen / Composites: Part A 33 (2002) 1209–1218 1217
18 A. Yasmin, P. Bowen/ Composites: Part A 33(2002)1209-1218 composite under quasi-static loads using a real-time [28] Mah T, Mendiratta MG, Katz AP, Mazdiyasni KS Recent develop- 正AP technique. J Compos Mater 1995; 29(13): 1680-94 cture and fatigue of a continuous fibre reinforced glass ceraMic composite. PhD Thesis. University of Birmingham, [291 Mah T, Mendiratta MG, Katz AP, Ruh R, Mazdiyasni KS. High- temperature mechanical behaviour of fibre-reinforced ceramic- 121 Kahraman R, Mandell JF. Influence of fibre-matrix bond strength e matrix composites. J Am Ceram Soc 1985: 68(9): C248-51 nance of Nicalon/CAS-ll composite J Co [30] Blissett M, Smith PA, Yeomans JA Flexural mechanical properties [22] Kerans RJ, Hay RS, Pagano NJ, Parthasarathy TA. The role of the calcium aluminosilicate composites. J Mater Sci 1998: 33: 4181-90 bre- reinforced interface in ceramic composites. Ceram Bull 1989: [31] Ishikawa T. Recent developments of the SiC fibre Nicalon and its 68(2):429-42. omposites, including properties of the SiC fibre Hi-Nicalon for ultra- [23] Kerans R, Parthasarathy TA Crack deflection in ceramic composites high temperature Compos Sci Technol 1994: 51: 135-44 and fibre coating design criteria. Composites, Part A 1999: 30:521 [32] Prewo KM, Johnson B, Starrett S. Silicon carbide fibre reinforced [24 Grande DH, Mandell JF, Hong KCC. Fibre-matrix bond strength glass ceramic composite tensile behaviour at elevated temperature udies of glass-ceramic and metal matrix composites. J Mater J Mater Sci1989;24:1373-9 88:23:311-28 [33] Yasmin A. Characterisation and evaluation of high temperatur [25] Davidge RW. Mechanical behaviour of ceramics. Cambridge solid composites. PhD Thesis. University of Birmingham, UK; 2001 tate science series. Cambridge University Press: 1979. [34] Elizalde MR, Sanchez JM, Daniel AM, Puente l, Martini A, Martinez. [26 Puente J, Sanchez JM, Elizalde R, Martinez JM, Fuentes M, Martin A Eznaola JM, Fuentes M. Damage evolution in ceramic matrix Effect of temperature on the tensile behaviour of CAS/SIC CMC ECF composites ECF 1996: 11: 1763-8. 1996;11:1703-8 [35] Daniel IM, Ishai O. Engineering mechanics of composite materials. [27] Allen RF, Percival MCL, Bowen P. Effects of test Oxford: Oxford University Press: 1994. loading rate on the fatigue and fracture resistance of a continuous fibre [36] Daniel IM, Anastassopoulos G. Failure mechanics and damage reinforced glass ceramic matrix composite. Proc Ninth Int Conf evolution in crossply ceramic-matrix composites. Int J Solids Struct Compos Mater 1993: 2: 121-8. 1995:32(314)341-55
composite under quasi-static loads using a real-time acoustic– ultrasonic NDE technique. J Compos Mater 1995;29(13):1680–94. [20] Allen RF. Fracture and fatigue of a continuous fibre reinforced glass ceramic matrix composite. PhD Thesis. University of Birmingham, UK; 1995. [21] Kahraman R, Mandell JF. Influence of fibre–matrix bond strength on the performance of Nicalon/CAS-II composite. J Compos Mater 1996; 30(8):864–84. [22] Kerans RJ, Hay RS, Pagano NJ, Parthasarathy TA. The role of the fibre-reinforced interface in ceramic composites. Ceram Bull 1989; 68(2):429–42. [23] Kerans RJ, Parthasarathy TA. Crack deflection in ceramic composites and fibre coating design criteria. Composites, Part A 1999;30:521–4. [24] Grande DH, Mandell JF, Hong KCC. Fibre–matrix bond strength studies of glass–ceramic and metal matrix composites. J Mater Sci 1988;23:311–28. [25] Davidge RW. Mechanical behaviour of ceramics. Cambridge solid state science series, Cambridge University Press; 1979. [26] Puente J, Sanchez JM, Elizalde R, Martinez JM, Fuentes M, Martin A. Effect of temperature on the tensile behaviour of CAS/SiC CMC. ECF 1996;11:1703–8. [27] Allen RF, Percival MCL, Bowen P. Effects of test temperature and loading rate on the fatigue and fracture resistance of a continuous fibre reinforced glass ceramic matrix composite. Proc Ninth Int Conf Compos Mater 1993;2:121–8. [28] Mah T, Mendiratta MG, Katz AP, Mazdiyasni KS. Recent developments in fibre-reinforced high temperature ceramic composites. Ceram Bull 1987;66(2):304–8. [29] Mah T, Mendiratta MG, Katz AP, Ruh R, Mazdiyasni KS. Hightemperature mechanical behaviour of fibre-reinforced ceramic– matrix composites. J Am Ceram Soc 1985;68(9):C248–51. [30] Blissett MJ, Smith PA, Yeomans JA. Flexural mechanical properties of thermally treated unidirectional and cross-ply Nicalon-reinforced calcium aluminosilicate composites. J Mater Sci 1998;33:4181–90. [31] Ishikawa T. Recent developments of the SiC fibre Nicalon and its composites, including properties of the SiC fibre Hi-Nicalon for ultrahigh temperature. Compos Sci Technol 1994;51:135–44. [32] Prewo KM, Johnson B, Starrett S. Silicon carbide fibre reinforced glass ceramic composite tensile behaviour at elevated temperature. J Mater Sci 1989;24:1373–9. [33] Yasmin A. Characterisation and evaluation of high temperature composites. PhD Thesis. University of Birmingham, UK; 2001. [34] Elizalde MR, Sanchez JM, Daniel AM, Puente I, Martini A, MartinezEznaola JM, Fuentes M. Damage evolution in ceramic matrix composites. ECF 1996;11:1763–8. [35] Daniel IM, Ishai O. Engineering mechanics of composite materials. Oxford: Oxford University Press; 1994. [36] Daniel IM, Anastassopoulos G. Failure mechanics and damage evolution in crossply ceramic–matrix composites. Int J Solids Struct 1995;32(3/4):341–55. 1218 A. Yasmin, P. Bowen / Composites: Part A 33 (2002) 1209–1218