Journal of the European Ceramic Society 17(1997)1485-1500 997 Elsevier Science Limite Printed in great Britain. All PII:S0955-2219097)00004-6 0955-221997s1 Carbon-Fibre-Reinforced (YMAS) Glass-Ceramic Matrix Composites. I Preparation, Structure and Fracture Strength Valerie Bianchi, Paul Goursat. Wharton Sinkler b Marc Monthioux erik menessierc aLMcts, URA CNRS 320, Faculte des Sciences, 123, Avenue Albert Thomas, 87060 Limoges Cedex, france bCEMES, UPRESA 6015 CNRS, B P. 4347, 31055 Toulouse Cedex, France Ceramique et Composites, B P. 7, 65460 Bazet, france (Received 9 July 1996, revised version received 29 July 1996, accepted 2 December 1996) Abstract 1 Introduction Unidirectional continuous carbon-fibre-reinforced During the last decades, the development of the glass-ceramic matrix composites are fabricated for aerospace industries has promoted research on sliding applications. The microstructure of the new materials that combine lightness and matrix has been characterized by x-ray diffraction favourable thermomechanical properties. Mond and the microtexture and structure of the fibres lithic ceramics meet these requirements and are (pitch-based and PAN-based) have been studied by able to endure high temperatures and corrosive transmission electronic microscopy. The different atmospheres but their applications are limited by fracture behaviours of the composites are described their brittleness and their low toughness. It has three-point bend test and fracture surface observa- been demonstrated that continuous fibres in cera tion by scanning electronic microscopy ) Due to the mic matrix composites(CMCs) prevent catas- thermal expansion mismatch between the fibres and trophic failure by giving them non-linear the matrix, a network of microcracks appears in the mechanical behaviour. Carbon-fibre-reinforced composites on cooling after hot-pressing. The glass-matrix composites have been studied for a relationship between the microcrack spacing and long time because of their very interesting thermo- the fracture behaviour suggests modifications of mechanical properties. -3 Subsequently, the use of the fibre-matrix bond. c 1997 Elsevier Science SiC fibre for strengthening a ceramic matrix has Limited allowed applications in air at moderate tempera tures to be considered. In particular, SiC-fibre- Des composites unidirectionnels a fibres longues de reinforced YMAS matrix composites have been carbone et matrice vitroceramique sont elabores proved to be very strong with a controlled frac pour des applications en frottement sec. La ture. Similarly, monolithic ceramics could be microstructre de la matrice est suivie par diffraction used for wear applications but again their brittle des rayons X. La micro-texture et la microstructure ness and their friction coefficient restrict their des fibres (ex-brai et ex-PAN) sont etudiees en development. For this reason researchers are more Microscopie electronique en Transmission. Les and more concerned with the use of solid lubri- differents comportement a la rupture des compo- cants and ceramic composites in which carbon sites sont ensuite decrits (essais en flexion trois fibres could contribute their lubricant properties points et observation des fractures en Microscopie However, many other factors must be considered Electronique a Balayage). Du fait des differ- and are of great importance, such as the geometry ences de dilatation thermique entre les fibres et of the contact between the mechanical parts,the la matrice, un reseau de fissures apparait au formation and the evolution of a third body from refroidissement dans les composites lors de leur mechanical or physicochemical rcactions at the frottage sous charge. La relation entre le pas de interface fissuration et le comportement d la rupture des This paper, which concerns the fabrication of composites suggere une evolution de la liaison unidirectional continuous carbon-fibre reinforced fibre-matrice YMAs glass-ceramic matrix composites for d 1485
JournalofrheEuropean CeramicSocieIy 17 (1997)1485-1500 0 1997 Elsevier Science Limited Printed in Great Britain. All rights reserved PII: SO955-2219(97)00004-6 0!355-22191971$17.OLl Carbon-Fibre-Reinforced (YMAS) Glass-Ceramic Matrix Composites. I. Preparation, Structure and Fracture Strength Val&ie Bianchi,” Paul Goursat,’ Wharton Sinkler,’ Marc Monthiouxb & Erik Mbnessier” “LMCTS, URA CNRS 320, Faculte des Sciences, 123, Avenue Albert Thomas, 87060 Limoges Cedex, France ‘CEMES UPRESA 6015 CNRS, B.P. 4347, 31055 Toulouse Cedex, France ‘Ceramiques et Composites, 1B.P. 7, 65460 Bazet, France (Received 9 July 1996, revised version received 29 July 1996, accepted 2 December 1996) Abstract 1 Introduction Unidirectional continuous carbon-jbre-reinforced glass-ceramic matrix composites are fabricated for dry sliding applications. The microstructure of the matrix has been characterized by X-ray dtJ?raction and the microtexture and structure of the fibres (pitch-based and PAN-based) have been studied by transmission electronic microscopy. The direrent fracture behaviours of the composites are described (three-point bend test and “fracture surface observation by scanning electronic microscopy). Due to the thermal expansion mismatch between the fibres and the matrix, a network of microcracks appears in the composites on cooling after hot-pressing. The relationship between the ,microcrack spacing and the fracture behaviour suggests modtjications of the jbre-matrix bond. 0 1997 Elsevier Science Limited. Des composites unidirectionnels a fibres longues de carbone et matrice vitroceramique sont elabores pour des applications en frottement sec. La microstructre de la matrice est suivie par dtjraction des rayons X. La micro-texture et la microstructure des fibres (ex-brai et ex-PAN) sont dtudiees en Microscopic Electronique en Transmission. Les d@rents comportements a la rupture des composites sont ensuite decrits (essais en flexion trois points et observation des fractures en Mcroscopie Electronique 6 Balayage). Du fait des d@rences de dilatation thermique entre les Jibres et la matrice, un reseau de fissures apparait au refroidissement dans les composites lors de leur frittage sous charge. La relation entre le pas de fissuration et le comportement h la rupture des composites suggere une &volution de la liaison fibre-matrice. During the last decades, the development of the aerospace industries has promoted research on new materials that combine lightness and favourable thermomechanical properties. Monolithic ceramics meet these requirements and are able to endure high temperatures and corrosive atmospheres but their applications are limited by their brittleness and their low toughness. It has been demonstrated that continuous fibres in ceramic matrix composites (CMCs) prevent catastrophic failure by giving them non-linear mechanical behaviour. Carbon-fibre-reinforced glass-matrix composites have been studied for a long time because of their very interesting thermomechanical properties. l-3 Subsequently, the use of SIC fibre for strengthening a ceramic matrix has allowed applications in air at moderate temperatures to be considered. In particular, SiC-fibrereinforced YMAS matrix composites have been proved to be very strong with a controlled fracture.4 Similarly, monolithic ceramics could be used for wear applications but again their brittleness and their friction coefficient restrict their development. For this reason researchers are more and more concerned with the use of solid lubricants and ceramic composites in which carbon fibres could contribute their lubricant properties. However, many other factors must be considered and are of great importance, such as the geometry of the contact between the mechanical parts, the formation and the evolution of a third body from mechanical or physicochemical reactions at the interface. This paper, which concerns the fabrication of unidirectional continuous carbon-fibre reinforced YMAS glass-ceramic matrix composites for dry 1485
chi et al sliding uses, is essentially devoted to the study of TEM), using a 120 kv high voltage and a Super the carbon fibres as affected by sintering to the twin objective stage. The microscope was also microstructural changes of the YMAS (Y, O3, equipped with energy dispersive analysis of the MgO, A1 O,, Sio,)matrix and to the mechanical X-ray system(EDAX), which was used to perform of the materials according to the ther- local and comparative chemical analysis on areas mal treatments larger than 50 nm and for elements with Z >5 2.1.2 Preparation of samples prior to investigations 2 Materials and Experimental Procedure As-received carbon fibres were prepared either by thin-sectioning of single filaments using a micro- 2.1 Carbon fibres tome(LKB Ultrotome) equipped with a diamond Three grades of fibre differing with respect to pre- knife(MS Type from DIATOME)or ion-thinning cursor and their final properties, pitch-based fibres of fibre tows according to the procedure described (Thornel P25 and P55)and a PAN-based fibre hereafter. Prior to both the preparation methods, (Torayca T400H) were used. Some properties of samples were embedded in an epoxy resin(TAAB these fibres are reported in table 1 or POlarbeD 812 any carbon fibre is highly anisotropic, with the trimmed then mechanically thing e tows were It is well known that the overall texture for Small pieces of resin-embedded fib usIng dia- graphenes (1 polyaromatic layer graphene) mond tools, then diamond and Sic powder oriented parallel to the fibre axis. However, the suspensions and abrasive papers. The final step of parallelism may be bctter or worse depending mechanical thinning was performed on both on the fabrication process and specifically on sample faces using a dimple grinder (GATAN parameters such as the precursor and/or the final 656)with a diamond powder suspension (grain carbonization temperature. In contrast, the textural sizes were in the 24 um range), in order to reach aspects of fibres may be much more variable thicknesses close to about 30 um. Thereafter, ion when seen on cross-sections, also depending on thinning was performed(GaTAN Duomill 600) the fabrication process, and specifically on param- using a 4 kv voltage and incidence angles from 20 eters such as the precursor (pitch or PAN, pitch to 12 grade,.)and/or the spinneret geometry. In addi tion, changes in texture may also be found within 2. 1. 3 Results a given fibre grade because of changes in the 2. 1.3.1 P55 fibres. The P55 fibres present aPan fabrication process that the manufacturer may Am'texture similar to that already found for have decided to introduce at some time without P55S and P75 AMOCO fibres 0 with internal changing the trade name. This may explain, for decohesion between polyaromatic sheets which instance, why cross-sections of T300 fibre from certainly originate in the structural shrinkage Torayca are nowadays found to exhibit some occurring during the final carbonization step, anisotropic features, , while this was never reported during which superimposed graphene layers al in the 1980s. Accurate characterization of the tend to move closer together with incrcasing ransversal texture of fibres used as reinforcement perature in any carbon materials(the lower limit is for the composites is therefore important that of the interlayer distance for graphite 0 335 nm- which is, however, practically never 2. 1. 1 Investigation methods reached in commercial carbon fibres) The micro- and nanotextural (isotropy, anis Electron microscopy allows textural hetero- otropy, porosity,.) and structural(crystallinity, geneities to be revealed. Specifically, large defects..study was mainly performed with a anisotropic areas (i.e made of graphenes roughly transmission electron microscope(Philips CM12 riented in a common direction at a long distance exhibit numerous and small( 100 nm) inclusions Table 1. Characteristics of fibres(given by the manufacturers) of isotropic (microporous)areas (i.e. made of graphenes with no preferred orientation, arrow in Fig. 1). In addition to the textural difference, the P25 P55 T400lI graphene stacks in the isotropic areas are less structurally organized, as shown by the lacl Longitudinal tensile strength(MPa) 1400 1900 450 Longitudinal tensile modulus(GPa) 160 380 250 ragg fringes(Bragg fringes originate from Tensile fracture elongation (% order Bragg reflections which are stopped by the Density (g/cm) obiective aperture an creforc makc wcll-ordered Filament diameter (um) Number of fibres per tow 20003000 phases appear locally darker in bright field imaging mode). As a matter of fact, higher magnifications
1486 V. Bianchi et al. sliding uses,5 is essentially devoted to the study of TEM), using a 120 kV high voltage and a Superthe carbon fibres as affected by sintering, to the twin objective stage. The microscope was also microstructural changes of the YMAS (Y,OJ, equipped with energy dispersive analysis of the MgO, Al,O,, SiO& matrix and to the mechanical X-ray system (EDAX), which was used to perform behaviours of the materials according to the ther- local and comparative chemical analysis on areas mal treatments. larger than 50 nm and for elements with Z > 5. 2 Materials and Experimental Procedure 2.1 Carbon fibres Three grades of fibre differing with respect to precursor and their final properties, pitch-based fibres (Thornel P25 and P55) and a PAN-based fibre (Torayca T400H) were used. Some properties of these fibres are reported in Table 1. 2.1.2 Preparation of samples prior to investigations As-received carbon fibres were prepared either by thin-sectioning of single filaments using a microtome (LKB Ultrotome) equipped with a diamond knife (MS Type from DIATOME) or ion-thinning of fibre tows according to the procedure described hereafter. Prior to both the preparation methods, samples were embedded in an epoxy resin (TAAB or POLARBED 812). It is well known that the overall texture for any carbon fibre is highly anisotropic, with the graphenes (1 polyaromatic layer = graphene) oriented parallel to the fibre axis. However, the parallelism may be better or worse depending on the fabrication process and specifically on parameters such as the precursor and/or the final carbonization temperature. In contrast, the textural aspects of fibres may be much more variable when seen on cross-sections, also depending on the fabrication process, and specifically on parameters such as the precursor (pitch or PAN, pitch grade, . ..) and/or the spinneret geometry. In addition, changes in texture may also be found within a given fibre grade because of changes in the fabrication process that the manufacturer may have decided to introduce at some time without changing the trade name. This may explain, for instance, why cross-sections of T300 fibre from Torayca are nowadays found to exhibit some anisotropic features,6,7 while this was never reported in the 198Os.* Accurate characterization of the transversal texture of fibres used as reinforcement for the composites is therefore important. Small pieces of resin-embedded fibre tows were trimmed then mechanically thinned using diamond tools, then diamond and Sic powder suspensions and abrasive papers. The final step of mechanical thinning was performed on both sample faces using a dimple grinder (GATAN 656) with a diamond powder suspension (grain sizes were in the 2-4 pm range), in order to reach thicknesses close to about 30 pm. Thereafter, ionthinning was performed (GATAN Duomill 600) using a 4 kV voltage and incidence angles from 20 to 12”. 2.1.3 Results 2.1.1 Investigation methods 2.1.3.1 P.55 j&es. The P55 fibres present a ‘PanAm’ texture9 similar to that already found for P55S and P75 AMOCO fibres’O with internal decohesions between polyaromatic sheets which certainly originate in the structural shrinkage occurring during the final carbonization step, during which superimposed graphene layers always tend to move closer together with increasing temperature in any carbon materials (the lower limit is that of the interlayer distance for graphite - 0.335 nm - which is, however, practically never reached in commercial carbon fibres). The micro- and nanotextural (isotropy, anisotropy, porosity, . ..) and structural (crystallinity, defects , . ..) study was mainly performed with a transmission electron microscope (Philips CM 12 Table 1. Characteristics of fibres (given by the manufacturers) Fibre P25 P55 T4OOH Longitudinal tensile strength (MPa) 1400 1900 4500 Longitudinal tensile modulus (GPa) 160 380 250 Tensile fracture elongation (%) 0.9 0.5 1.8 Density (gkm3) 1.90 2.0 1.80 Filament diameter (pm) 11 10 7 Number of fibres per tow 2000 2000 3000 Electron microscopy allows textural heterogeneities to be revealed. Specifically, large anisotropic areas (i.e. made of graphenes roughly oriented in a common direction at a long distance) exhibit numerous and small (= 100 nm) inclusions of isotropic (microporous) areas (i.e. made of graphenes with no preferred orientation, arrow in Fig. 1). In addition to the textural difference, the graphene stacks in the isotropic areas are less structurally organized, as shown by the lack of Bragg fringes (Bragg fringes originate from high order Bragg reflections which are stopped by the objective aperture and therefore make well-ordered phases appear locally darker in bright field imaging mode). As a matter of fact, higher magnifications
Carbon-fibre-reinforced(YMAS) glass-ceramic matrix composites. I 1487 chemical heterogeneities within the pitch precu sor. 2 As a general rule for the thermal behaviour of polyaromatic carbon phases, the microporous part will always stay far from the structural state of genuine graphite, whatever the heat-treatment tcmpcraturc. In contrast, the anisotropic parts are more likely to evolve towards the graphite struc- ture(without ever reaching it, however), and are esponsible for the relatively high modulus value and low ultimate strength and strain of the P55 fbre relative to the PAN-based T400H fibre (Table 1). Finally, TEM observations( Figs 2 and 3)show that the directions of preferred orien tations may vary greatly at a scale of a hundred nanometers and lower which is not in the resolu tion range for optical microscopy(about 0.5 um) Hence, the Pan-Am texture revealed microscopy is only an average statistical tendency At the fibre surface a third texture often occurs It is made of very stiff (i.e. free of misorientations) graphene sheets whicl an incomplete coating around the fibre. The structural organ 100mm ization is very good: N may be 50 or more and la may reach 100 nm(Fig 3). Such a carbon skin is ig. 1. P55 fibre cross-section. Bright field TEM image. White not systematic and does not adhere well to the areas within anisotropic(dark) areas are decohesion between fibre either polyaromatic sheets. The grey part (arrow) is an isotropic inclusion 2.1.3. 2 P25 fibres. Low magnification TEM images show that the fibre surface may appear either (Fig. 2)illustrate the structural difference between smooth (Fig. 4(a)) or corrugated (Fig 4(b)) the isotropic(above) and the anisotropic(below) Generally speaking, the textural and chemical fea- regions. In as much as the crystalline order is tures of the P25 fibres are similar to those of P55 biperiodic only (in contrast to genuine graphite), fibres when observed by TEM, i.e. a combination the structural state of such polyaromatic carbon of isotropic and anisotropic areas and a highly materials, so-called turbostratic, may be charac- anisotropic but incomplete polyaromatic carbon terized by the average lateral extension La of skin. However, the structures are similar but actu graphenes and the average number N of graphenes ally not identical. Indeed, the structural organiz- stacked N and La thus define the average poly- ation of graphene stacks is lower, as clearly aromatic entity which may be considered coherent revealed by high resolution imaging(Fig. 5). The relative to the diffraction electron beam. Regard- number n of graphenes stacked into coherent ing the structural difference between the isotropic entities is in the 5-10 range and their extension L and the anisotropic parts in the fibre, La and n (within anisotropic domains) is most often 2-3 nm are 3-5 nm and 4-8, respectively. for the former, (as compared to 40-50 and 10 nm, respectively while N may reach 40-50 and La may reach 10 nm for the P55 fibre). A consequence is that grain for the latter. Each entity is somewhat misoriented boundaries are more numerous than in P55 fibres relative to the neighbouring entities, and separated since the coherent entities are smaller, which from them by grain boundaries. In addition, induces a higher ultimate strength valuc(Tablc 1) Fig 2 illustrates some details of Bragg fringes Indeed, the higher the number of boundaries, the (arrows)and also how decohesion may occur higher the number of linking sites preventing the within anisotropic areas (clear zones between easy gliding between two superimposed graphenes natic sheets). An important feature, demon- For the same reason, the bulk modulus of the p25 strated by investigations on longitudinal sections fibre is lower than the P55 fibre f fibres(not illustrated), is that the microporous Finally, the features are all consistent with th (isotropic) areas are not spherical but cylindrical fact that P25 fibres originate from the same pre- inclusions within the fibre, more or less flattened, cursor and have undergone the same fabrication with the cylinder axis being parallel to the fibre process as P55 fibres, but with a lower final axis. Such a textural heterogeneity is due to carbonization temperature This explains why the
Carbon-Jbre-reinforced ( YMAS) glass-ceramic matrix composites. I, 1487 Fig. 1. P55 fibre cross-section. Bright field TEM image. White areas within anisotropic (dark) aireas are decohesions between polyaromatic sheets. The grey part (arrow) is an isotropic inclusion. (Fig. 2) illustrate the structural difference between the isotropic (above) and the anisotropic (below) regions. In as much as the crystalline order is biperiodic only (in contrast to genuine graphite), the structural state of such polyaromatic carbon materials, so-called turbostratic, may be characterized by the average lateral extension L, of graphenes and the average number N of graphenes stacked. N and L, thus dlefine the average polyaromatic entity which may be considered coherent relative to the diffraction electron beam. Regarding the structural difference between the isotropic and the anisotropic parts in the fibre, L, and N are 3-5 nm and 48, respectively, for the former, while N may reach 40-50 and L, may reach 10 nm for the latter. Each entity is somewhat misoriented relative to the neighbouring entities, and separated from them by grain boundaries.” In addition, Fig.2 illustrates some details of Bragg fringes (arrows) and also how decohesions may occur within anisotropic areas (clear zones between aromatic sheets). An important feature, demonstrated by investigations on longitudinal sections of fibres (not illustrated), is that the microporous (isotropic) areas are not spherical but cylindrical inclusions within the fibre., more or less flattened, with the cylinder axis being parallel to the fibre axis. Such a textural heterogeneity is due to chemical heterogeneities within the pitch precursor. I2 As a general rule for the thermal behaviour of polyaromatic carbon phases, the microporous part will always stay far from the structural state of genuine graphite, whatever the heat-treatment temperature. In contrast, the anisotropic parts are more likely to evolve towards the graphite structure (without ever reaching it, however), and are responsible for the relatively high modulus value and low ultimate strength and strain of the P55 fibre relative to the PAN-based T400H fibre (Table 1). Finally, TEM observations (Figs 2 and 3) show that the directions of preferred orientations may vary greatly at a scale of a hundred nanometers and lower, which is not in the resolution range for optical microscopy (about 0.5 pm). Hence, the Pan-Am texture revealed by optical microscopy’ is only an average statistical tendency. At the fibre surface, a third texture often occurs. It is made of very stiff (i.e. free of misorientations) graphene sheets which act as an incomplete coating around the fibre. The structural organization is very good: N may be 50 or more and L, may reach 100 nm (Fig.3). Such a carbon skin is not systematic and does not adhere well to the fibre either. 2.1.3.2 P25fib res. Low magnification TEM images show that the fibre surface may appear either smooth (Fig. 4(a)) or corrugated (Fig 4(b)). Generally speaking, the textural and chemical features of the P25 fibres are similar to those of P55 fibres when observed by TEM, i.e. a combination of isotropic and anisotropic areas and a highly anisotropic but incomplete polyaromatic carbon skin. However, the structures are similar but actually not identical. Indeed, the structural organization of graphene stacks is lower, as clearly revealed by high resolution imaging (Fig. 5). The number N of graphenes stacked into coherent entities is in the 5-10 range and their extension L, (within anisotropic domains) is most often 2-3 nm (as compared to 40-50 and 10 nm, respectively, for the P55 fibre). A consequence is that grain boundaries are more numerous than in P55 fibres, since the coherent entities are smaller, which induces a higher ultimate strength value (Table 1). Indeed, the higher the number of boundaries, the higher the number of linking sites preventing the easy gliding between two superimposed graphenes. For the same reason, the bulk modulus of the P25 fibre is lower than the P55 fibre Finally, the features are all consistent with the fact that P25 fibres originate from the same precursor and have undergone the same fabrication process as P55 fibres, but with a lower final carbonization temperature. This explains why the
1488 y Bianchi et al interlayer decohesion, usual in P55 fibres(Fig. 2), ring evidenced at about half or third of the radius have not occurred in P25 fibres(Fig. 5). since ( Fig. 6(a) ). This feature is consistent with previous the structural shrinkage has not been intense observations performed on the T300 grade of PAN-based fibres from Torayca which exhibit a coarsely isotropic texture with a local anisotropic 2.1.3.3 T400H fibres. The ovcrall molphology of radial texture in the ring.6,7A possible explanation the cross-sections may be round, oval or bean. is that this is an effect of the mechanical stretching shaped, with diameters in the 6-8 un range and a imposed on the green fibres during the spinning ISOTROPIC INCLUSION ANISOTROPIC AREA.'i 10 nm Fig. 2. P55 fibre cross-section. Lattice fringe TEM image. Above is an enlarged view of an isotropic area. Below is an enlarged
1488 V. Bianchi et al. interlayer decohesions, usual in P55 fibres (Fig. 2), have not occurred in P25 fibres (Fig. 5), since the structural shrinkage has not been intense enough. 2.2.3.3 T400H fibres. The overall molphology of the cross-sections may be round, oval or beanshaped, with diameters in the 6-8 pm range and a ring evidenced at about half or third of the radius (Fig. 6(a)). This feature is consistent with previous observations performed on the T300 grade of PAN-based fibres from Torayca which exhibit a coarsely isotropic texture with a local anisotropic radial texture in the ring.6,7 A possible explanation is that this is an effect of the mechanical stretching imposed on the green fibres during the spinning Fig. 2. P55 fibre cross-section. Lattice fringe TEM image. Above is an enlarged view of an isotropic area. Below is an enlarged view of an anisotropic area. Arrows illustrate Bragg fringes
Carbon-fibre-reinforced(YMAS) glass-ceramic matrix composites. I 1489 step, which may induce internal compressive/ten- originate from some outgassing event urring sile stress discrimination, in association with the within the fibre when still softenable. at the first radial oxygen gradient developed during the steps of the fibering process. Furthermore, subsequent curing and carbonization step. Based observation is again similar to previous investi on non-convincing TEM observations, the same tions on T300 fibres. 3 or T400H fibres which, in authors have proposcd a twinned radial/concentric addition, have demonstrated?9 that the pores are texture for the dark ring in T300 fibres. This is consistently cylindrical instead of spherical unlikely for several reasons. Among others, their axis oriented parallel to the fibre axis. High revealing a radial orientation resolution imaging(Fig. 8)also reveals that the high resolution imaging(see below), and never a only anisotropic domain is encountered within the concentric one ring(which appears dark in low magnification Low magnification TEM images provide infor- images: Fig. 6(a) where the overall orientation of mation on the fibre surface roughness(Figs 6(a) graphenes is perpendicular to the fibre surface and 6(b). The surface is rougher than that of the (radial texture). In addition, the graphene stacks P25 and P55 fibres, and is consistent with the pre- have a better structural organization than within vious investigations performed on T300 fibres. the surrounding isotropic part of the fibre High resolution TEM imaging allows the nano- The textural and structural features of the textural and structural aspects to be described T400H fibre explain the mechanical properties (Fig. 7). The structural organization is very poor, Table 1). For the same reasons as those given for the These features are close to that of P25 fibres, but the highest, while the elastic modulus is low th is since N is about 3-7 and La is about 1-3 nm. P55/P25 fibre comparison, the ultimate stren of long range anisotropy (in cross-section). Another specific fea- 2.2 YMAS glass- ceramic matrix ture is the occurrence of round pores(arrow The glass powder is obtained by Fig. 7)which are found in any T400H fibre, pref- (Y2O3, MgO, AlO3, SiO2), melting, cooling and erentially in areas close to the fibre surface grinding to obtain a grain size lower than 10 um arrows in Fig. 6(b). They pre-exist within the The content of each oxide is chosen to yield, after fibre, i.e. they are not induced by the preparation the thermal treatment of the glass, a glass-ceramic procedure for the TEM investigations, for instance, mainly consisting of cordierite, yttrium silicates since they are not found in other ion-thinned and a small quantity of the MgAl,O4 spinel phase PAN-based fibres from another source but are The glass transition(T and crystallization(To quite similar in other respects. Such holes probably temperature intervals have been determined by 10m Fig-3 P25 or P55 fibre cross-section. Lattice fringe TEM image. Example of the very well-structured polyaromatic carbon phase found at the fibre surface
Carbon-Jibre-reinforced ( YMAS) glass-ceramic matrix composites. I. 1489 step, which may induce internal compressive/tensile stress discrimination, in association with the radial oxygen gradient developed during the subsequent curing and carbonization step.13 Based on non-convincing TEM observations, the same authors have proposed a twinned radial/concentric texture for the dark ring in T300 fibres. This is unlikely for several reasons. Among others, we succeeded in revealing a radial orientation in high resolution imaging (see below), and never a concentric one. Low magnification TEM images provide information on the fibre surfa.ce roughness (Figs 6(a) and 6(b)). The surface is rougher than that of the P25 and P55 fibres, and is consistent with the previous investigations performed on T300 fibres.7*13 High resolution TEM imaging allows the nanotextural and structural aspects to be described (Fig. 7). The structural organization is very poor, since iV is about 3-7 and L, is about l-3 nm. These features are close to that of P25 fibres, but a major difference is the lack of long range anisotropy (in cross-section). Another specific feature is the occurrence of round pores (arrow in Fig. 7) which are found in any T400H fibre, preferentially in areas close to the fibre surface (arrows in Fig. 6(b)). They pre-exist within the fibre, i.e. they are not induced by the preparation procedure for the TEM investigations, for instance, since they are not found in other ion-thinned PAN-based fibres from another source but are quite similar in other respects.’ Such holes probably originate from some outgassing event occurring within the fibre when still softenable, at the first steps of the fibering process. Furthermore, this observation is again similar to previous investigations on T300 fibres7,13 or T400H fibres’ which, in addition, have demonstrated739 that the pores are consistently cylindrical instead of spherical, with their axis oriented parallel to the fibre axis. High resolution imaging (Fig. 8) also reveals that the only anisotropic domain is encountered within the ring (which appears dark in low magnification images: Fig.6(a)) where the overall orientation of graphenes is perpendicular to the fibre surface (radial texture). In addition, the graphene stacks have a better structural organization than within the surrounding isotropic part of the fibre. The textural and structural features of the T400H fibre explain the mechanical properties (Table 1). For the same reasons as those given for the P55/P25 fibre comparison, the ultimate strength is the highest, while the elastic modulus is low. 2.2 YMAS glass-ceramic matrix The glass powder is obtained by mixing oxides (Y,03, MgO, A1203, SiO& melting, cooling and grinding to obtain a grain size lower than 10 pm. The content of each oxide is chosen to yield, after the thermal treatment of the glass, a glass-ceramic mainly ,consisting of cordierite, yttrium silicates and a small quantity of the MgAl*O, spine1 phase. The glass transition (T,) and crystallization (T,) temperature intervals have been determined by Fig. 3. P25 or P55 fibre cross-section. Lattice fringe TEM image. Example of the very well-structured polyaromatic carbon phase found at the fibre surface
1490 V. Bianchi et al differential thermal analysis(dta)of the glass It can be noted that the cte of the glass powder(Setaram Micro ATD M5--linear increase ceramic is lower than that of the glass, which of5°Cmin),as815855°Cand945-1055° depends on the low Cte of the cordierite. The respectively glass transition temperature interval obtained with The coefficients of thermal expansion(CTE) of this technique is similar to that determined by the glass and the glass-ccramic have bcc DTA sured with a vertical dilatometer(Setaram TMA 2), in argon, with a linear heating or cooling rate 2.3 Compusites processing (C/min)on bars obtained by unidirectional Pre-preg sheets were prepared by using a slurry pressing and sintering in air for 2h at the required infiltration process of fibre tows. The impregnated temperature. The results are summarized in Table fibres are wound on a hexagonal mandrel. The volume fraction of fibres is nearly equal to 35%. After drying, the tapes were cut into 35 x 35 mm pieces, stacked and debound. Then the samples were densified in argon by hot-pressing 950-1250oC temperature interval with a pressure of 10 MPa(LPA-DVM Goliath) followed by linear increase or decrease in temperature (25.C/min). The unidirectional pressure was applied when the sintering temperature had been reached and was removed on cooling In order to optimize the density and the fracture strength of the composites, various thermal cycles have been performed. The conditions are summarized in Table 3 where in some cases. two temperature-time combinations are given. The first corresponds to the densification step under plied pressure. The second -- if any sponds to an additional pressureless crystallization step at a higher temperature. H山r 3 Composites Microstructure and Properties 3.1 Density and fibl The density has been measured by the hydrostatic technique For the P25 fibre-reinforced composites, after hot-pressing for at least 30 min, the open porosity ranges from 1. 5 to 3. 7%0, except for Pl, which was hot-pressed at a low temperature and the open porosity of which is about 24%. Indeed, at 950oC, the viscosity of the glass is too high to allow a vis cous flow in the tows Nevertheless. some bundles of fibres were not impregnated by the glass and the distribution of the fibres in the matrix is not homogeneous(Fig 9(a) The density of T400H fibre-reinforced composites increases gradually as the sintering tempcrature is FBER raised up to 1150C and the viscosity decreases bably due to the high ber of fibres tow(3000), the composites always exhibit some porosity (Table 4) and non-impregnated bundles Fig. 4. P25 fibre cross-sections. Bright field TEM image, of fibres can be pulled out during polishing taken from fibres within the same composite(P5). The matrix Fig 9(b). For sintering temperatures higher than appears dark. Clear areas within the fibres are isotropic parts while dark areas are anisotropic parts: (a)and (b) are two 1150oC, the porosity increases again. This latter examples of fibres with various corrugated surfaces. phenomenon could be related to a release of nitro
1490 V. Biunchi et al. differential thermal analysis (DTA) of the glass powder (Setaram Micro ATD MS-linear increase of SWmin), as 815-855°C and 945-1055”C, respectively. The coefficients of thermal expansion (CTE) of the glass and the glass-ceramic have been measured with a vertical dilatometer (Setaram TMA 92), in argon, with a linear heating or cooling rate (3Wmin) on bars obtained by unidirectional pressing and sintering in air for 2h at the required temperature. The results are summarized in Table 2 5.14 Fig. 4. P25 fibre cross-sections. Bright field TEM image, taken from fibres within the same composite (PS). The matrix appears dark. Clear areas within the fibres are isotropic parts, while dark areas are anisotropic parts; (a) and (b) are two examples of fibres with various corrugated surfaces. It can be noted that the CTE of the glassceramic is lower than that of the glass, which depends on the low CTE of the cordierite. The glass transition temperature interval obtained with this technique is similar to that determined by DTA. 2.3 Composites processing Pre-preg sheets were prepared by using a slurry infiltration process of fibre tows. The impregnated fibres are wound on a hexagonal mandrel. The volume fraction of fibres is nearly equal to 35%. After drying, the tapes were cut into 35 X 35 mm’ pieces, stacked and debound. Then the samples were densified in argon by hot-pressing in the 950-1250°C temperature interval with a pressure of 10 MPa (LPA-DVM Goliath) followed by a linear increase or decrease in temperature (25”Umin). The unidirectional pressure was applied when the sintering temperature had been reached and was removed on cooling. In order to optimize the density and the fracture strength of the composites, various thermal cycles have been performed. The conditions are summarized in Table 3 where, in some cases, two temperature-time combinations are given. The first corresponds to the densification step under applied pressure. The second - if any - corresponds to an additional pressureless crystallization step at a higher temperature. 3 Composites Microstructure and Properties 3.1 Density and fibre distribution The density has been measured by the hydrostatic technique. For the P25 fibre-reinforced composites, after hot-pressing for at least 30 min, the open porosity ranges from 1.5 to 3.7%, except for Pl, which was hot-pressed at a low temperature and the open porosity of which is about 24%. Indeed, at 95O”C, the viscosity of the glass is too high to allow a viscous flow in the tows. Nevertheless, some bundles of fibres were not impregnated by the glass and the distribution of the fibres in the matrix is not homogeneous (Fig. 9(a)). The density of T4OOH fibre-reinforced composites increases gradually as the sintering temperature is raised up to 1150°C and the viscosity decreases. Probably due to the high number of fibres in a tow (3000), the composites always exhibit some porosity (Table 4) and non-impregnated bundles of fibres can be pulled out during polishing (Fig. 9(b)). For sintering temperatures higher than 115O”C, the porosity increases again. This latter phenomenon could be related to a release of nitro-
Carbon-fibre-reinforced (YMAS) glass-ceramic matrix composites. I gen from T400H fibres. Such a release occurring the viscosity of the matrix may become too high when the matrix is still viscous (at 1050 for T7) because of crystallization, which could prevent the likely to prevent good adhesion between the fibres impregnation of the fibres and the matrix whereas, for hot-pressing at 1150C (T5), carbon pyrolysis is nearly achieved 3. 2 Matrix crystallization le matrix can closely take the forms of the X-ray diffraction equipment with a computer fibres. When hot-pressing occurs at 1200C (T6), search/match program was used to determine the ANISOTROPIC AREA 310m Fig. 5. P25 fibre cross-section, Lattice fringe TEM image. The comparison with Fig. 2 demonstrates the lower structural organization f P25 fibres relative to P55 fibres, while exhibiting similar textural features (anisotropic domains). Note the absence of interlayer decohesion
Carbon-jbre-reinforced ( YMAS) glass-ceramic matrix composites. I. 1491 gen from T400H fibres. Such a release occurring when the matrix is still vis’cous (at 1050” for T7) is likely to prevent good adhlesion between the fibres and the matrix wherea,s, for hot-pressing at 1150°C (T5), carbon pyrolysis is nearly achieved and the matrix can closely take the forms of the fibres. When hot-pressing occurs at 1200°C (T6), the viscosity of the matrix may become too high because of crystallization, which could prevent the impregnation of the fibres. 3.2 Matrix crystallization X-ray diffraction equipment with a computer search/match program was used to determine the Fig. 5. P25 fibre cross-section. Lattice fringe TEM image. The comparison with Fig. 2 demonstrates the lower structural organization of P25 fibres relative to P55 fibres, while exhibiting similar textural features (anisotropic domains). Note the absence of interlayer decohesions
1492 chi et al nature of the phases and to follow the crystalliza- tion of the matrix versus the sintering conditions The XRd results were obtained at room tempera ture from the sintered samples Whatever the fibres used. the microstructural changes in the matrix were similar(Fig. 10(a, b)) Up to 1000C, X-ray diffraction patterns show that the matrix remains mainly vitreous. Only the diffraction lines of a-alumina appear, but this compound is also detected in the starting glass powder. The MgAl,O4 spinel (JCPDS 22-1152 file) precipitates above 1050C, as well as the hexagonal high temperature cordierite (JCPDS 13-0293 file indialite) and the a-and B-Y2Si2O, yttrium ilicates (JCPDs 38-0223, 38-0440 files). The SiO2-Y,O, system is very complex with many polymorphic phases in the case of Y?,O,(a, B, Y, 8. The phase transformations depend on temp- rature, time and impurities. The crystalliza starts around 1050c but the rate is slow and after one hour it seems to stop At 1250.C, the a-Y2S120, turns almost completely into the B form. This result agrees with the transformation temperature indicated by Liddel and Thompson. 5 ccurring within the matrix with increasing 1 un m temperature is being published as a companion 3.3 Microcracking and fracture strength of the obvious for their coefficient of thermal expansion which is close to zero along the fibre axis and about 20 x 106C. in the radial direction. 7 The thermal expansion mismatch with the matrix induces stresses on cooling. When these residual stresses are higher than the matrix strength, a net- work of microcracks, which are perpendicular to the fibre axis and are deflected by the fibres, can form in the matrix and are observed in the present composites lacocca and Duquettehave observed the same microcracking phenomenon in carbon-fibre-reinforced glass-matrix composites and have studied their effects on the oxidation behaviour of comp sites. Different models have been proposed to determine the interfacial stresses9-25 in order to compare them with the mechanical strength of the matrix. Observation of micro- racking on metallographically prepared samples showed that the cracks were regularly spaced. The average distance between neighbouring cracl Fig. 6. T400H fibre cross-section. Bright field TEM image changed with the hot-pressing conditions of the taken from fibres within a composite(T5).(a) Note the sizes composites(Table 5), which means that the inter and shapes of fibres, together with the dark ring at about one facial shear friction stress can change with the well-corrugated fibre surface. arrows indicate hole located sintering temperature. The A C K. theory allows close to the fibre surface
1492 V. Bianchi et al. nature of the phases and to follow the crystallization of the matrix versus the sintering conditions. The XRD results were obtained at room temperature from the sintered samples. Whatever the fibres used, the microstructural changes in the matrix were similar (Fig. lO(a, b)). Up to lOOO”C, X-ray diffraction patterns show that the matrix remains mainly vitreous. Only the diffraction lines of a-alumina appear, but this compound is also detected in the starting glass powder. The MgAl,O, spine1 (JCPDS 22-1152 file) precipitates above lOSO”C, as well as the hexagonal high temperature cordierlite (JCPDS 13-0293 file - indialite) and the a-and p-Y&O7 yttrium silicates (JCPDS 38-0223, 38-0440 files). The SiO*-Y20, system is very complex with many polymorphic phases in the case of Y,Si,07 (a, p, y, 8). The phase transformations depend on temperature, time and impurities. The crystallization starts around 1050°C but the rate is slow and after one hour it seems to stop. At 125O”C, the cw-Y,Si,07 turns almost completely into the p form. This result agrees with the transformation temperature indicated by Liddel and Thompson.” An extensive study of the structural changes occurring within the matrix with increasing temperature is being published as a companion paper. ” 3.3 Microcracking and fracture strength of the composites The anisotropic structure of the carbon fibre used leads to anisotropic properties. In particular this is obvious for their coefficient of thermal expansion which is close to zero along the fibre axis and about 20 x 10m6 ‘C-l in the radial direction.17 The thermal expansion mismatch with the matrix induces stresses on cooling. When these residual stresses are higher than the matrix strength, a network of microcracks, which are perpendicular to the fibre axis and are deflected by the fibres, can form in the matrix and are observed in the present composites. Iacocca and Duquette” have observed the same microcracking phenomenon in carbon-fibre-reinforced glass-matrix composites and have studied their effects on the oxidation behaviour of composites. Different models have been proposed to determine the interfacial stresses’9-25 in order to compare them with the mechanical strength of the matrix. Observation of microcracking on metallographically prepared samples showed that the cracks were regularly spaced. The average distance between neighbouring cracks changed with the hot-pressing conditions of the composites (Table 5), which means that the interfacial shear friction stress can change with the sintering temperature. The A.C.K. theory” allows Fig. 6. T400H fibre cross-section. Bright field TEM images taken from fibres within a composite (T5).(a) Note the sizes and shapes of fibres, together with the dark ring at about one third of the fibre radius. (b) Enlargement revealing the WellLcorrugated fibre surface. Arrows indicate hole located close to the fibre surface
Carbon fibre-reinforced(YMAS) glass-ceramic matrix composites. I 1493 the interfacial shear stress T to be estimated from can be pointed out that the crack spacing changes the distance I between microcracks less significantly for P25 fibre-reinforced compo sites than for T400H fibre-reinforced composites (1) That would indicate that, after debonding the relative sliding movement between fibre and where Vm and Vf are the matrix and fibre volume matrix would dissipate more or less energy. fractions, respectively, r is the radius of the fibre Indeed, the nature and the resistance to debond and omu is the fracture strength of the matrix. It ing and sliding of the fibre-matrix interface can 10 nm Fig. 7. T400H fibre cross-section. Bright field TEM image taken in the light-grey part of a fibre from Fig. 6. The presence of holes (arrow) attests that the area imaged is close to the fibre surface
Carbon-Jibre-reinforced ( YMAS) glass-ceramic matrix composites. I. 1493 the interfacial shear stress T to be estimated from can be pointed out that the crack spacing changes the distance I between microcracks : less significantly for P25 fibre-reinforced compo- +W sites than for T400H fibre-reinforced composites. 2.7 (1) That would indicate that, after debonding, the relative sliding movement between fibre and where V, and V, are the matrix and fibre volume matrix would dissipate more or less energy. fractions, respectively, r is the radius of the fibre Indeed, the nature and the resistance to debondand a,, is the fracture strength of the matrix. It ing and sliding of the fibre-matrix interface can Fig. 7. T4OOH fibre cross-section. Bright field TEM image taken in the light-grey part of a fibre from Fig. 6. The presence of holes (arrow) attests that the area imaged is close to the fibre surface
proceed from physicochemical reactions at the have different structures. 6 Moreover, their CTE interface and from the thermal stresses induced can also be different and non-linear with tempera y the differences between the expansion coeffi- inducing different stresses in the matrix. he components. It is possible that Other parameters like the radius and the Youngs physicochemical reactions between the P25 fibre modulus of the fibre influence the development occur between the T400H fibre and the matrix, For example, with a higher fibre young 0.? and the matrix arc diffcrcnt from thosc which of thermal stresses and matrix cracking strain since the fibres come from different precursors and lus, the stress in the matrix increases. Indeed Fig 8. T400H fibre cross-section. Bright field TEM image taken in the dark-grey ring within the fibre from Fig. 6. The distribution of the graphene stacks is highly anisotropic( they are radially oriented relative to the fibre surface)
1494 V. Bianchi et al. proceed from physicochemical reactions at the interface and from the thermal stresses induced by the differences between the expansion coefficients of the components. It is possible that physicochemical reactions between the P25 fibre and the matrix are different from those which occur between the T400H fibre and the matrix, since the fibres come from different precursors and have different structures.*‘j Moreover, their CTE can also be different and non-linear with temperature,i4’ ” inducing different stresses in the matrix. Other parameters like the radius and the Young’s modulus of the fibre influence the development of thermal stresses and matrix cracking strain. For example, with a higher fibre Young’s modulus, the stress in the matrix increases. Indeed, Fig. 8. T4OOH fibre cross-section. Bright field TEM image taken in the dark-grey ring within the fibre from Fig. 6. The distribution of the graphene stacks is highly anisotropic (they are radially oriented relative to the fibre surface)
