Acta mater. vol 44 Copyright C 1996 Acta Metallurgica 09567151(95)00386X Printed in Great Dr 1359645496515.00+0.00 MICROSTRUCTURE-PROPERTY RELATIONSHIPS OF SIC FIBRE-REINFORCED MAGNESIUM ALUMINOSILICATES-I. MICROSTRUCTURAL CHIARACTERISATION A KUMAR+ and K M. KNOWLES University of Cambridge, Department of Materials Science and Metallurgy, Pembroke Street, Cambridge CB2 3QZ, England (Received 31 October 1994: in revised form 6 September 1995) Abstract-The microstructure of two magnesium aluminosilicates unidirectionally reinforced with SiC fibres(Nicalon) has been examined. a diphasic interlayer having a higher O/Si ratio than in the fibres was found on the surface of the fibres in both composites. This interlayer could be identified as an amorphous mixture of silica and carbon in the composite hot-pressed just below the liquidus temperature of stoichiometric cordierite(composite A). In the other composite hot-pressed at 920 C and subsequently craned at 1150C (composite B), a relatively thicker diphasic interlayer was observed, consisting of lentified in composite A. A thin interlayer consisting mostly of matrix elements was also identified etween the diphasic interlayer and the discrete carbon interlayer in this cor le. Dierences in structure and morphology of interfacial regions in the two composites could clearly be attributed differences in the hot-pressing schedules. The basal planes of turbostratic carbon were aligned parallel he fibre-matrix interfaces in both composites. Copyright c 1996 Acia Metallurgica Inc. 1 INTRODUCTION It is evident from the early work of Brennan [1, 2 Glass-ceramics reinforced with continuous Sic fibres offer a combination of properties which are desirable related to the nature of interfacial layers or high-temperature structural applications. the an [1, 2] demonstrated that a carbon-rich layer high strength and high modulus of these composites at the fibre matrix interface is always present in is a direct consequence of the high strength and tough SiC/lithium aluminosilicate(LAs)composites modulus of the nanocrystalline SiC fibres, while the investigate the nature of the number of researchers have since attempted to damage tolerance of these composites arises from the ese interfacial layers and ability of fibres first to carry load after transverse ve tried to correlate this with the mechanical pull out from the matrix properties of other potential glass-ceramic matrix materials. For these toughening mechanisms to oper- aluminosilicates (MAS). Chaim et al. B3 observed a relatively weak in shear to allow debonding at these Sic/MAS composite which they interpreted as silica interfaces, yet strong enough to give good load because of the brittle nature of the composite. They transfer between the fibre and the matrix. Therefore. were unable to identify any carbon-rich layers unam the mechanical properties of these composites, and in oguously, even in the damage tolerant composites Metcalfe et al. [4] have recentl particular their damage tolerance, depend largely on demonstrated that strong and tough composites with the mechanical properties of the fibre-matrix inter faces, which in turn are controlled by the structure MAS as the matrix can be produced, but they did not analyse the matrix-fibre interfaces in sufficient and chemistry of the interfaces. Hence, in order to detail to be able to identify the structure and chem- there is a need to characterise both the chemistry and istry of the thin reaction layer they. observed by the microstructure of the fibre-matrix interface and scanning electron microscopy. It is, therefore, the phases present in the matrix apparent that an unambiguous identification of the structure and the chemistry of the interfacial layers Present address: Department of Mechanical mate performance of Naval Postgraduate School, Montere U.S.A literature
Pergamon 0956-7151(95)00386-X Acta mater. Vol. 44, No. 7, PP. 2901-2921. 1996 Copyright 0 1996 Acta Metallurgica Inc. Published by Elsevier Science Ltd Printed in Great Britain. All rights reserved 1359-6454/96 $15.00 + 0.00 MICROSTRUCTURE-PROPERTY RELATIONSHIPS OF Sic FIBRE-REINFORCED MAGNESIUM ALUMINOSILICATES-I. MICROSTRUCTURAL CHARACTERISATION A. KUMARt and K. M. KNOWLES University of Cambridge, Department of Materials Science and Metallurgy, Pembroke Street, Cambridge CB2 342, England (Received 31 October 1994; in revised form 6 September 1995) Abstract-The microstructure of two magnesium aluminosilicates unidirectionally reinforced with Sic fibres (Nicalon) has been examined. A diphasic interlayer having a higher OjSi ratio than in the fibres was found on the surface of the fibres in both composites. This interlayer could be identified as an amorphous mixture of silica and carbon in the composite hot-pressed just below the liquidus temperature of stoichiometric cordierite (composite A). In the other composite hot-pressed at 920°C and subsequently ceramed at 1150°C (composite B), a relatively thicker diphasic interlayer was observed, consisting of turbostratic carbon together with amorphous silica. A distinct interlayer of turbostratic carbon was identified in composite A. A thin interlayer consisting mostly of matrix elements was also identified between the diphasic interlayer and the discrete carbon interlayer in this composite. Differences in the structure and morphology of interfacial regions in the two composites could clearly be attributed to differences in the hot-pressing schedules. The basal planes of turbostratic carbon were aligned parallel to the fibre-matrix interfaces in both composites. Copyright 0 1996 Acfa Metallurgica Inc. 1. INTRODUCTION Glass-ceramics reinforced with continuous SIC fibres offer a combination of properties which are desirable for high-temperature structural applications. The high strength and high modulus of these composites is a direct consequence of the high strength and modulus of the nanocrystalline SIC fibres, while the damage tolerance of these composites arises from the ability of fibres first to carry load after transverse matrix cracking and then to pull out from the matrix beyond the peak in the load-deflection curve for such materials. For these toughening mechanisms to operate successfully, the fibre-matrix interfaces must be relatively weak in shear to allow debonding at these interfaces, yet strong enough to give good load transfer between the fibre and the matrix. Therefore, the mechanical properties of these composites, and in particular their damage tolerance, depend largely on the mechanical properties of the fibre-matrix interfaces, which in turn are controlled by the structure and chemistry of the interfaces. Hence, in order to optimise the mechanical properties of composites, there is a need to characterise both the chemistry and the microstructure of the fibre-matrix interface and the phases present in the matrix. TPresent address: Department of Mechanical Engineering, Naval Postgraduate School, Monterey. CA 93943, U.S.A. It is evident from the early work of Brennan [l, 21 that the mechanical properties of composites are closely related to the nature of interfacial layers. Brennan [l, 21 demonstrated that a carbon-rich layer at the fibre-matrix interface is always present in tough Sic/lithium aluminosilicate (LAS) composites. A number of researchers have since attempted to investigate the nature of these interfacial layers and have tried to correlate this with the mechanical properties of other potential glass-ceramic matrix composites, such as BC fibre-reinforced magnesium aluminosilicates (MAS). Chaim et al. [3] observed a reaction layer at matrix-fibre interfaces in a SiC/MAS composite which they interpreted as silica because of the brittle nature of the composite. They were unable to identify any carbon-rich layers unambiguously, even in the damage tolerant composites they produced. Metcalfe et al. [4] have recently demonstrated that strong and tough composites with MAS as the matrix can be produced, but they did not analyse the matrix-fibre interfaces in sufficient detail to be able to identify the structure and chemistry of the thin reaction layer they observed by scanning electron microscopy. It is, therefore, apparent that an unambiguous identification of the structure and the chemistry of the interfacial layers and their effects on the ultimate performance of SiC/MAS composites has not yet been reported in the literature
KUMAR and KNOWLES: SiC REINFORCED ALUMINOSILICAtESI In the work reported here, two different mag- sitional maps. Electron beam energy and beam cur- nesium aluminosilicate matrices unidirectionally rent were 5kv and 4nA, respectively. A digital reinforced with SiC (Nicalon)-fibres have been exam- resolution of 128 x 128 was used to collect all the ned to determine the relationship between mechan- compositional images using a conventional eds de ical properties and microstructure. In this paper(Part tector (detection limit atomic number >11).Dwell D), the evolution of microstructure during processing time per pixel was 2 s for composite A and 3 s for and post-processing heat-treatments is described for composite B No matrix correction was applied to the both composites. Qualitative and quantitative analy- digitised maps and therefore these maps are used only ses of interfacial layers and the microstructure of the for comparison purpose composites have been carried out using conventional 2.2. Interface analysis. The fibre-matrix inter- and analytical electron microscopy, optical faces were analysed using thin foils in transmission microscopy and X-ray diffraction techniques. The electron microscopes. Both a Philips 400T operating mechanical behaviour of these composites is reported at 120 kV and a JEOL 2000FX operating at 200kV in Part I [S were used. A JEOL 4000EX-lI operating at 400 kv was used to perform high-resolution electron mi- 2. MATERIALS AND EXPERIMENTAL METHODS croscopy(HREM). An ultra thin window energy dispersive X-ray analyser(EDS, Link model 6284) and a parallel electron energy-loss spectrometer Composite plates of sizes 10 x 10 cm in the form of (PEELS, Gatan model 666) on the JEOL 2000FX layers of unidirectional fibre-reinforced matrices, were used to carry out the compositional and struc. ea[0l, configuration, were received from Pilkington tural analysis of the fibres, the matrices and th Technology Centre, Lathom, U.K. in two batches. fibre-matrix interfaces SiC(Nicalon, NL 202, Nippon Carbon Company Thin foils for TEM examination were prepared by of Japan)fibres were used in both batches. A mechanically thinning the small rectangular samples glass of stoichiometric cordierite composition to thicknesses of A 20-50 um. The mechanical thin (2MgO2Al2O, 5SiO2)was chosen as the matrix ning was carried out only on high grade SiC papers material for composite A, whereas a glass of compo- (600 and 800)and the pressure applied was kept to weight with small amounts of P2 Os(2.0 wt%)and time than the routine grinding process, it reduced the B,O, (1.0 wt%)was used for composite B. P2O, acts erosion of fibre-matrix interfacial areas during mech anical thinning. Both surfaces of the mechanically B,O, lowers the melt temperature and delays particle surface crystallisation [6, 7]. As will be discussed thinned samples were polished for approx 15 min using first 6 um diamond paste and then l um Section 3. 1, the rationale for choosing this compo- diamond paste. This was essential for uniform thin sition is to reduce the hot-pressing temperature ning of samples in an ion-beam thinner. The polished glass for both composites was melted in air in a central hole of size 2000 or 1000 u u using sil gas-fired furnace. The composite plates were made by the slurry impregnation procedure for making con- as adhesive. This was followed by ion-beam tinuous fibre-reinforced composites [8]. The plates ( Gatan dual ion mill model 600) at 5kv were hot-pressed inert atmosphere at 1500 and incidence to perforation, and finally thinning at a low 920C for composites A and B, respectively 9 Plates of composite B were ceramed in air for I h at 1150C after consolidation, whereas composite a did not siC(220) receive any heat-treatment after hot-pressing SiC (11) 2. 2. Experimental c(110) 2.2. 1. Microstructural characterisation, The micro- structure of both composites was assessed using an X-ray diffractometer(Philips PW 3719), an optical microscope (Olympus, model BHM) and scanning electron microscopes (SEM, Camscan S2 and Radial Camscan $4). Polished composite cross-sections were coated with either carbon or gold before SEM exam ination. Energy dispersive X-ray analysis(EDS)of fibre and matrix was carried out on a sem(camscan S4)using a conventional Link EDS analyser with a detection limit of atomic number >11 Fig. 1 Schematic SADP illustrating overlapping rings from a dedicated clectron-probe microanalyser Sio,, C and Sic along with the aperture positions for radial ( Cameca, SX-50) was used to obtain digital compo- oration of the diffraction pattern
2902 KUMAR and KNOWLES: Sic REINFORCED ALUMINOSILICATES-I In the work reported here, two different magnesium aluminosilicate matrices unidirectionally reinforced with SIC (Nicalon)-fibres have been examined to determine the relationship between mechanical properties and microstructure. In this paper (Part I), the evolution of microstructure during processing and post-processing heat-treatments is described for both composites. Qualitative and quantitative analyses of interfacial layers and the microstructure of the composites have been carried out using conventional and analytical electron microscopy, optical microscopy and X-ray diffraction techniques. The mechanical behaviour of these composites is reported in Part II [5]. 2. MATERIALS AND EXPERIMENTAL METHODS 2.1. Materials Composite plates of sizes 10 x 10 cm in the form of n layers of unidirectional fibre-reinforced matrices, i.e. a [0], configuration, were received from Pilkington Technology Centre, Lathom, U.K. in two batches. SIC (Nicalon, NL 202, Nippon Carbon Company of Japan) fibres were used in both batches. A glass of stoichiometric cordierite composition (2Mg0.2A1203.5Si02) was chosen as the matrix material for composite A, whereas a glass of composition 22.3% MgO, 21.3% Al,O,, 53.4% SiO, by weight with small amounts of P,O, (2.0 wt%) and B,03 (1.0 wt%) was used for composite B. P,O, acts as a nucleating agent and reduces viscosity, whereas B,O, lowers the melt temperature and delays particle surface crystallisation [6,7]. As will be discussed in Section 3.1, the rationale for choosing this composition is to reduce the hot-pressing temperature required to consolidate the composites. The precursor glass for both composites was melted in air in a gas-fired furnace. The composite plates were made by the slurry impregnation procedure for making continuous fibre-reinforced composites [8]. The plates were hot-pressed in an inert atmosphere at 1500 and 920°C for composites A and B, respectively [9]. Plates of composite B were ceramed in air for 1 h at 1150°C after consolidation, whereas composite A did not receive any heat-treatment after hot-pressing. 2.2. Experimental 2.2.1. Microstructural characterisation. The microstructure of both composites was assessed using an X-ray diffractometer (Philips PW 3719), an optical microscope (Olympus, model BHM) and scanning electron microscopes (SEM, Camscan S2 and Camscan S4). Polished composite cross-sections were coated with either carbon or gold before SEM examination. Energy dispersive X-ray analysis (EDS) of fibre and matrix was carried out on a SEM (Camscan, S4) using a conventional Link EDS analyser with a detection limit of atomic number 2 11. A dedicated electron-probe microanalyser (Cameca, SX-50) was used to obtain digital compositional maps. Electron beam energy and beam current were 5 kV and 4nA, respectively. A digital resolution of 128 x 128 was used to collect all the compositional images using a conventional EDS detector (detection limit atomic number > 11). Dwell time per pixel was 2 s for composite A and 3 s for composite B. No matrix correction was applied to the digitised maps and therefore these maps are used only for comparison purposes. 2.2.2. Interface analysis. The fibre-matrix interfaces were analysed using thin foils in transmission electron microscopes. Both a Philips 400T operating at 120 kV and a JEOL 2000FX operating at 200 kV were used. A JEOL 4000EX-II operating at 400 kV was used to perform high-resolution electron microscopy (HREM). An ultra thin window energy dispersive X-ray analyser (EDS, Link model 6284) and a parallel electron energy-loss spectrometer (PEELS, Gatan model 666) on the JEOL 2000FX were used to carry out the compositional and structural analysis of the fibres, the matrices and the fibre-matrix interfaces. Thin foils for TEM examination were prepared by mechanically thinning the small rectangular samples to thicknesses of x20-50 pm. The mechanical thinning was carried out only on high grade SIC papers (600 and 800) and the pressure applied was kept to a minimum. Although this procedure took a longer time than the routine grinding process, it reduced the erosion of fibre-matrix interfacial areas during mechanical thinning. Both surfaces of the mechanically thinned samples were polished for approximately 15 min using first 6 pm diamond paste and then 1 pm diamond paste. This was essential for uniform thinning of samples in an ion-beam thinner. The polished sections were then mounted onto copper rings with a central hole of size 2000 or 1000 pm using silver paint as adhesive. This was followed by ion-beam thinning (Gatan dual ion mill model 600) at 5 kV and 15” incidence to perforation, and finally thinning at a low silica Fig. 1. Schematic SADP illustrating overlapping rings from SiO,, C and Sic along with the aperture positions for radial exploration of the diffraction pattern
KUMAR and KNOWLEs: SIC REINFOKCED ALUMINOSILICA'TES-I 2%U3 angle(a10)for 5-30 min to extend the thin area. Table 1. Crystal structure and lattice parameters of the phases When required, the thin foils for hREM were coated with a thin layer of amorphous carbon using a carbon Pha Structure evaporator unit(Edwards E306) Quantitative and qualitative analyses of a micro- a-cordieriteb structural feature require a detailed correlation Orthorhombic 9.2 between the diffraction pattern and the image. The 'A.S.T. M index card no. 15-776 centred dark field technique was therefore used to a.s.t.m index card no, 11-273 dentify the phases present at interfacial layers. This vas achieved by radial exploration of reciprocal space following the procedure described by Oberlin was lower than that which could be readily detected [10 by the diffractometer(realistically an amount <10 A schematic diagram depicting the relative pos- wt %, even if a sophisticated analysis of the x-ray pattern(SADP), from SiO2, SiC and C is shown in were to have been used). Cordierite was present in the Fig.I.Amorphous silica gives a diffuse ring at high temperature hexagonal form in both composites 0.41 nm and a faint plateau out to 0.12 nm[11]. The The crystal structure and lattice parameters of the B-SiC in the nanocrystalline Nicalon fibres gives 111 relevant phases are given in Table 1. It was also noted 251 nm), 220(0. 154 nm)and 311(0.131 nm)rings. that in composite A the diffraction peaks of B-siC 02(0.344 nm), 100(0.212 nm)and 110(0.11 nm)(from the fibres)were better defined than those seen rings of turbostratic carbon are also shown in Fig. 1. in composite B. This suggests that in composite A the Turbostratic carbon consists of aromatic layer stacks average grain size of crystals of p-Sic in the fibres that are piled up in parallel but rotated slightly at Optical micrographs of polished cross-sections of ndom relative to one another [12 ] This structure is characterised by the presence of the 002 reflection composites A and B are shown in Figs 2(a)and(b), 44 A)and hko instead of hkI diffraction spots respectively. Qualitatively, the fibre distribution is or diffraction rings from graphite [12]. When the relatively uniform, but some matrix-rich regions are bjective aperture is in position 1(Fig. 1), C and Sio, also evident. Mean fibre volume fractions estimated will appear bright in a dark field image, whereas in using an optical microscope were 0. 47 and 0.40 for position 2(Fig. 1)crystals of Sic will appear bright. composites A and B, respectively. In general, the When imaging interfacial regions it is important polished sections of composite A are pore-free, that the interfacial layer is parallel to the electron whereas micro-pores are clearly seen on the polished beam.The specimen needs to be tilted so that the surfaces of composite B. Back-scattered electron terfacial layer is aligned parallel to the electron images of the polished sections of composites A and beam(edgc-on imaging). If the interface layer is not B are shown in Figs 3(a) and(b), respcctivcly. A seen edge-on, misinterpretation of the interface 1-2 um thick bright layer was consistently observed gions is possible because of overlapping of around the fibres in composite a [Fig 3(a)l, whereas fibre-interface-matrix regions (see for example a very thin dark layer was seen around the fibres in the discussion in Ref. [13]) composite B[Fig. 3(b)]. The different morphologies of the crystalline phases within the matrix are also 3. RESULTS AND INTERPRETATION readily seen in Figs 2 and 3 The observed differences in the microstructure of 3. 1. Microstructural characterisation the two composites arise from the difference in Analysis of X-ray diffraction (XRD) traces of processing temperature and the composition of glass compositesAandBrevealedthefollowingmulliteprecursorsusedtofabricatethecomposites.com (AlO3 2SiO2) was the major crystalline phase in posite A was hot-pressed at a temperature just below which small amounts of the liquidus temperature of stoichiometric cordierite a-cordierite(2MgO. 2A1,O3. 5SiO,), a-cristobalite glass. The formation of different phases on cooling (SiO,)and relatively large amounts of residual glass from the hot-pressing temperature can be understood were present, and (i)a-cordierite and protoenstatite by considering the pseudobinary section going (MgSiO,)were the major crystalline phases in com. through cordierite(2MgO. 2Al2O3. 5SiO2) and the osite B, although some diffraction peaks corre- theoretical compound "Mg-beryl" in the sponding to orthoenstatite were also present in the MgO-Al2Oj-Sio, ternary phase diagram(Fig. 4, XRD trace of composite B. However, it is difficult to [15D. At temperatures below 1523 C mullite and gla distinguish the different pyroxenes because of the (i.e. supercooled liquid) will be present (point A overlapping of various peaks in XRD. The absence of Fig. 4). On further cooling to 1465 C (point B in Fig a hump, characteristic of a siliceous glassy phase, 4)all glass and mullite should react to form cordier between 20 and 30 20 in the XRD trace of com- ite. The co-existence of mullite, glass and,most istobalite with cordierite at room ss or that, more likely, the amount of residual glass temperature in composite a would suggest that the
KUMAR and KNOWLES: Sic REINFORCED ALUMINOSILICATES-I 2903 angle (x 100) for 5-30 min to extend the thin area. When required, the thin foils for HREM were coated with a thin layer of amorphous carbon using a carbon evaporator unit (Edwards E306). Quantitative and qualitative analyses of a microstructural feature require a detailed correlation between the diffraction pattern and the image. The centred dark field technique was therefore used to identify the phases present at interfacial layers. This was achieved by radial exploration of reciprocal space following the procedure described by Oberlin [lOI. A schematic diagram depicting the relative positions of various hkl rings in a selected area diffraction pattern (SADP), from SiOz, SIC and C is shown in Fig. 1. Amorphous silica gives a diffuse ring at 0.41 nm and a faint plateau out to 0.12 nm [ll]. The /?-Sic in the nanocrystalline Nicalon fibres gives 111 (0.251 nm), 220 (0.154 nm) and 311 (0.131 nm) rings. 002 (0.344 nm), 100 (0.212 nm) and 110 (0.11 nm) rings of turbostratic carbon are also shown in Fig. 1. Turbostratic carbon consists of aromatic layer stacks that are piled up in parallel but rotated slightly at random relative to one another [12]. This structure is characterised by the presence of the 002 reflection (d > 3.44 A) and hk0 instead of hkl diffraction spots or diffraction rings from graphite [12]. When the objective aperture is in position 1 (Fig. 1), C and SiO, will appear bright in a dark field image, whereas in position 2 (Fig. 1) crystals of SIC will appear bright. When imaging interfacial regions it is important that the interfacial layer is parallel to the electron beam. The specimen needs to be tilted so that the interfacial layer is aligned parallel to the electron beam (edge-on imaging). If the interface layer is not seen edge-on, misinterpretation of the interface regions is possible because of overlapping of fibre-interface-matrix regions (see for example the discussion in Ref. [13]). 3. RESULTS AND INTERPRETATION 3.1. Microstructural characterisation Analysis of X-ray diffraction (XRD) traces of composites A and B revealed the following: (i) mullite (3Al,0,.2Si02) was the major crystalline phase in composite A, in addition to which small amounts of cr -cordierite (2MgO. 2Al,O,. 5SiO,), cr -cristobalite (SiO,) and relatively large amounts of residual glass were present, and (ii) c( -cordierite and protoenstatite (MgSiO,) were the major crystalline phases in composite B, although some diffraction peaks corresponding to orthoenstatite were also present in the XRD trace of composite B. However, it is difficult to distinguish the different pyroxenes because of the overlapping of various peaks in XRD. The absence of a hump, characteristic of a siliceous glassy phase, between 20” and 30” 28 in the XRD trace of composite B suggested that either there is no residual glass or that, more likely, the amount of residual glass Table 1. Crystal structure and lattice parameters of the phases present in the as-received composites Lattice parameters (A) Phase structure a b c Mullit? Orthorhombic 1.54 1.86 2.884 a-cordier@ Hexagonal 9.11 9.71 9.352 Protoenstatitec Orthorhombic 9.25 8.75 5.320 ‘A.S.T.M index card no. 15-716. bA.S.T.M index card no. 13-293. ‘A.S.T.M index card no. 1 l-273. was lower than that which could be readily detected by the diffractometer (realistically an amount < 10 wt%, even if a sophisticated analysis of the X-ray spectrum such as a modified Rietveld analysis [14] were to have been used). Cordierite was present in the high temperature hexagonal form in both composites. The crystal structure and lattice parameters of the relevant phases are given in Table 1. It was also noted that in composite A the diffraction peaks of B-Sic (from the fibres) were better defined than those seen in composite B. This suggests that in composite A the average grain size of crystals of /I-Sic in the fibres was higher than that in the fibres in composite B. Optical micrographs of polished cross-sections of composites A and B are shown in Figs 2(a) and (b), respectively. Qualitatively, the fibre distribution is relatively uniform, but some matrix-rich regions are also evident. Mean fibre volume fractions estimated using an optical microscope were 0.47 and 0.40 for composites A and B, respectively. In general, the polished sections of composite A are pore-free, whereas micro-pores are clearly seen on the polished surfaces of composite B. Back-scattered electron images of the polished sections of composites A and B are shown in Figs 3(a) and (b), respectively. A l-2 pm thick bright layer was consistently observed around the fibres in composite A [Fig. 3(a)], whereas a very thin dark layer was seen around the fibres in composite B [Fig. 3(b)]. The different morphologies of the crystalline phases within the matrix are also readily seen in Figs 2 and 3. The observed differences in the microstructure of the two composites arise from the difference in the processing temperature and the composition of glass precursors used to fabricate the composites. Composite A was hot-pressed at a temperature just below the liquidus temperature of stoichiometric cordierite glass. The formation of different phases on cooling from the hot-pressing temperature can be understood by considering the pseudobinary section going through cordierite (2MgO. 2A1, 03. 5Si02) and the theoretical compound “Mg-beryl” in the Mg@Al,03-Si02 ternary phase diagram (Fig. 4, [ 151). At temperatures below 1523°C mullite and glass (i.e. supercooled liquid) will be present (point A in Fig. 4). On further cooling to 1465°C (point B in Fig. 4) all glass and mullite should react to form cordierite. The co-existence of mullite, glass and, most importantly, cristobalite with cordierite at room temperature in composite A would suggest that the
2904 KUMAR and KNOWLES: SiC REINFORCED ALUMINOSILICATES-I 88 ● °° 50 um Fig. 2. Optical micrographs of polished cross-sections of (a) composite A and (b)composite B composition of the original glass precursor was non- mullite and liquid will be very sluggish and will stoichiometric and that it should be closer to the require very slow cooling rates. The two main reasons .2O-Sior side of the Mgo-AlO Sio, ternary for the co-existence of mullite, cordierite, cristobalite phase diagram(the glass composition would be in the and liquid in composite a are therefore: (i a small phase field on the extreme right of the pseudobinary deviation from the exact 2: 2: 5 cordierite compo- section in Fig. 4). This assumption is supported by sition and (i)a non-equilibrium cooling rate. he fact that the projection point of the 2: 2: 5 compo- Cristobalite and cordierite precipitate on further sition is completely surrounded by fields containing cooling. Cristobalite is undesirable because it under liquids as seen in an isothermal section at 1460.c goes a displacive phase transformation around Fig. 5, [15]. Therefore, if a glass composition closer 200-250'C to form low cristobalite [16]. This involves to the Al2O-SiO2 side of the ternary system is a substantial volume reduction (e 3.9%)which can chosen, mullite, cordierite and liquid would be pre- cause extensive cracking within cristobalite and sent at 1460'C. Furthermore, peritectic reactions are therefore in the matrix [17, 18 very slow and even in metallic systems they require A glass of non-stoichiometric cordierite compo very slow cooling rates to go to completion. There- sition in the MAs phase field was used to fabricate fore, the kinetics of the formation of cordierite from composite B. The two main reasons for selecting a
290 4 KUMAR and KNOWLES: Sic REINFORCED ALUMINOSILICATES-I F ?g. 2. Optical micrographs of polished cross-sections of (a) composite A and (b) composite B. composition of the original glass precursor was nonstoichiometric and that it should be closer to the A&O,-SiO, side of the Mg&Al,OrSiO, ternary phase diagram (the glass composition would be in the phase field on the extreme right of the pseudobinary section in Fig. 4). This assumption is supported by the fact that the projection point of the 2 : 2 : 5 composition is completely surrounded by fields containing liquids as seen in an isothermal section at 1460°C (Fig. 5, [ 151). Therefore, if a glass composition closer to the Al,03-SiO, side of the ternary system is chosen, mullite, cordierite and liquid would be present at 1460°C. Furthermore, peritectic reactions are very slow and even in metallic systems they require very slow cooling rates to go to completion. Therefore, the kinetics of the formation of cordierite from mullite and liquid will be very sluggish and will require very slow cooling rates. The two main reasons for the co-existence of mullite, cordierite, cristobalite and liquid in composite A are therefore: (i) a small deviation from the exact 2: 2: 5 cordierite composition and (ii) a non-equilibrium cooling rate. Cristobalite and cordierite precipitate on further cooling. Cristobalite is undesirable because it undergoes a displacive phase transformation around 200-250°C to form low cristobahte [ 161. This involves a substantial volume reduction (z 3.9%) which can cause extensive cracking within cristobalite and therefore in the matrix [17, 181. A glass of non-stoichiometric cordierite composition in the MAS phase field was used to fabricate composite B. The two main reasons for selecting a
KUMAR and KNOWLES: SIC REINFORCED ALUMINOSILICATES-I non-stoichiometric glass composition were: (i) to [20] proposed that cooling from the protoenstatite decrease the hot-pressing temperature and (ii)to stability field results in an intergrowth structure avoid the formation of mullite and cristobalite. The consisting of regions of clinoenstatite and orthoen composites were hot-pressed at 920.C to avoid crys- statite, the relative amounts of which will depend on tallisation during pressing, which would increase the the cooling rate Slower cooling rates are predicted to viscosity of the glass, Cordierite and protoenstatite produce more regions of orthoenstatite [20]. On the formed on ceraming the as-pressed composites at other hand, Schreyer and Schairer [15]used their own 1 150C. The presence of protoenstatite in the matrix experimental evidence to propose that protoenstatite at room temperature is surprising because this form could be retained at room temperature with very low of enstatite is known to be stable only at high amounts of clinoenstatite and orthoenstatite-like temperature [19-21]. Other forms of enstatite are: (i) phases when Al2O, was present in the bulk compo clinoenstatite, which forms on quenching proton- sition(see compositions no 35 and 37 in their paper) statite to room temperature, and which may be The formation of protoenstatite in similar MAS metastable at all temperatures; and (i)orthoenstatite, glasses has also been observed by many other re- a low temperature phase [19, 20]. Buseck and Iijima searchers, and yet none of them commented on the a b Fig. 3. Back-scattered electron images of polished cross-sections of (a)composite A and (b)
KUMAR and KNOWLES: Sic REINFORCED ALUMINOSILICATES-I 2905 non-stoichiometric glass composition were: (i) to decrease the hot-pressing temperature and (ii) to avoid the formation of mullite and cristobalite. The composites were hot-pressed at 920°C to avoid crystallisation during pressing, which would increase the viscosity of the glass. Cordierite and protoenstatite formed on ceraming the as-pressed composites at 1150°C. The presence of protoenstatite in the matrix at room temperature is surprising because this form of enstatite is known to be stable only at high temperature [19-211. Other forms of enstatite are: (i) clinoenstatite, which forms on quenching protoenstatite to room temperature, and which may be metastable at all temperatures; and (ii) orthoenstatite, a low temperature phase [19,20]. Buseck and Iijima [20] proposed that cooling from the protoenstatite stability field results in an intergrowth structure consisting of regions of clinoenstatite and orthoenstatite, the relative amounts of which will depend on the cooling rate. Slower cooling rates are predicted to produce more regions of orthoenstatite [20]. On the other hand, Schreyer and Schairer [ 151 used their own experimental evidence to propose that protoenstatite could be retained at room temperature with very low amounts of clinoenstatite and orthoenstatite-like phases when A&O, was present in the bulk composition (see compositions no. 35 and 37 in their paper). The formation of protoenstatite in similar MAS glasses has also been observed by many other researchers, and yet none of them commented on the Fig. 3. Back-scattered electron images of polished cross-sections of (a) composite A and (b) composite B