- Email: [email protected]
Microstructural investigation of interfaces in CMCs
Composites: Part A 33 (2002) 1467–1470 www.elsevier.com/locate/compositesa
Microstructural investigation of interfaces in CMCs G. Boitier*, S. Darzens, J.-L. Chermant, J. Vicens LERMAT, URA CNRS 1317, ISMRA, 6 Bd Mare´chal Juin, 14050 Caen Cedex, France
Abstract This paper is focused on the importance of pyrocarbon interfaces in two types of ceramic matrix composites (Cf – SiC and SiCf – SiBC), during creep tests under argon. The development of micromechanisms which consume energy and then allow a damage tolerance, depends on the morphology of the fiber/matrix interphase, which has been investigated by TEM and HRTEM. q 2002 Elsevier Science Ltd. All rights reserved. Keywords: A. Ceramic-matrix composites (CMCs); B. Interface/interphase; TEM; B. Creep
1. Introduction Ceramic matrix composites (CMCs) are a class of materials developed for aeronautics and space applications, in a domain where superalloys cannot be used anymore. They have potential applications in structures (air intakes, structural panels with stiffners, high dimensional stability structures for mirror or antenna, etc.) or in turbines (rear frame liners, mixer flow, petals, exhaust cones, etc.), or for brakes [1 – 4]. But their mechanical behavior depends mainly on the fiber/matrix interfaces (or interphases!). They can be characterized at the macroscopic level by debonding energy, G; and frictional interface shear stress, t; but understanding the different behaviors and obtained values requires investigations at the micro- and nanoscopic scale using TEM and HRTEM. The micro- and nanostructures can be correlated to the characteristics of the G and t values, and then to the life time parameters [5 –9]. The scope of this paper is to present the complexity, the role, the evolution and the importance of interfaces and interphases in two CMCs reinforced with continuous carbon or silicon carbide fibers: Cf –SiC and SiCf – SiBC, before and after creep tests.
2. Materials and techniques Cf –SiC and SiCf – SiBC were fabricated by Snecma Propulsion Solide (St Me´dard en Jalles, France) from chemical vapor infiltration (CVI) processes in woven 2.5 D fiber architectures (ex-PAN carbon fibers or Nicalon NLM 202 silicon carbide fibers). SiBC matrix is a self-healing * Corresponding author.
multilayered matrix, based on Si – C, B – C and Si – B –C phases. All the fiber architectures have received a thin pyrolytic carbon deposit before the infiltration of the matrix. Specimens were creep tested in tension, under a partial pressure of argon (500 mbar), in a temperature domain up to 1673 K and a stress domain between 110 and 250 MPa [10, 11]. Due to the mismatch of the coefficients of thermal expansion between fibers and matrix, Cf –SiC in the asreceived state has some matrix microcracks, which are not present in SiCf –SiBC composites.
3. Results During creep these composites exhibit a typical damagecreep [10 – 12], with matrix microcracking, fiber/matrix and yarn/matrix debonding, fiber and yarn bridging, fiber and yarn pull-out and rupture. So the deformation of these CMCs can be analyzed in terms of damage mechanics, as proposed by Kachanov [13]. The microcrack network will depend on the interaction with the pyrolitic carbon (PyC) deposited on the fibers, for Cf – SiC and SiCf –SiBC, and also with the matrix multilayers for SiCf –SiBC: the role of these interphases appears as predominant. The observation of the as-received composites reveals in both cases turbostratic carbon regions, globally parallel to the fibers and the matrix (Fig. 1), as already observed [14, 15] and classified by Despre`s [15] as ‘isotropic globally oriented’ in his classification of possible PyC interphases in Cf – SiC composites. In the case of Cf – SiC the observations and analysis of the PyC interphases reveal two types of interfacial sliding mechanisms [8 –10]. At 1473 K there is the interphase rupture by debonding between two carbon
1359-835X/02/$ - see front matter q 2002 Elsevier Science Ltd. All rights reserved. PII: S 1 3 5 9 - 8 3 5 X ( 0 2 ) 0 0 1 4 7 - 1
1468
G. Boitier et al. / Composites: Part A 33 (2002) 1467–1470
Fig. 1. HRTEM micrographs of the matrix/pyrocarbon interface in a as-received Cf – SiC (a) and of the matrix/pyrocarbon/fiber interfaces in a as-received SiCf –SiBC (b).
leaves: then the interfacial sliding can be assimilated to a dry friction between two rough solids. Close to the fibers and the matrix the structure of the PyC is turbostratic type, with atomic carbon planes parallel to the fiber direction. That has also clearly been evidenced by SEM observations of the surface of the fibers: that surface is very rough (see the small insert in Fig. 2a). Although at 1473 K the fiber/matrix interface is radially in compression, under stress some lenticular pores appear (Fig. 2a). These lenticular pores due to local decohesions between carbon planes are, in fact, the nuclei for the microcrack development. When the debonding is not total, carbon ribbons bridge the two parts of the microcracks. At 1673 K, the PyC interphase is a little degraded over about 100 nm from the matrix; there are a disappearing of the previous anisotropic texture, and a certain amorphization; then the interfacial sliding can be assimilated to a viscous flow (Fig. 2b). That was also confirmed by SEM observations: the fiber surfaces are smooth [8 –10] (see the small insert in Fig. 2b). If observations are performed at a higher scale one notes that the nanotexture of the carbon fibers increases with the test temperature. Consequently it leads to an increase of the local molecular orientation of oriented volumes of carbon planes parallel to the fibers. Although this phenomenon has been called ‘carbon fiber nanocreep’ [16], this contribution of the fibers and matrix to the macroscopic deformation is negligible, but it can be considered as the nuclei of the damage process. In the case of SiCf – SiBC same type of features are observed, but a little more complex due to the existence of the different matrix layers based on the Si –B – C system. In addition to the damage observed in Cf – SiC composites, there are also matrix microcrack deviations via some matrix multilayers where carbon exists [11]: this carbon is located at the interfaces of specific matrix layers and presents a thin
Fig. 2. HRTEM and SEM micrographs of the pyrocarbon interphase in Cf –SiC creep tested specimen under 220 MPa and in Ar: (a) with the presence of some lenticular pores at 1473 K; (b) and the presence of carbon rollings at 1673 K. Inserts correspond to SEM images of the surface features of the SiCf fibers at these two temperatures.
G. Boitier et al. / Composites: Part A 33 (2002) 1467–1470
1469
Fig. 3. TEM micrographs of the pyrocarbon interphase showing a microcrack deviation mode I ! mode II in a longitudinal section (parallel to the stress direction) (a), and a bridging of a microcrack by carbon ribbons (b), for SiCf –SiBC specimens creep tested at 1473 K, under 200 MPa and in Ar.
layer of PyC, from 10 to 100 nm thick, and oriented globally parallel to these interfaces. The larger the interphase is, the more pronounced the debonding between the matrix layers appears. In the case of the pyrocarbon layer (between the silicon carbide fibers and the matrix), close to the matrix the carbon structure is turbostratic type over a thickness of about 12 – 15 nm, i.e. 30– 35 carbon atomic planes. Further away, the carbon structure is globally isotropic oriented, according to the Despre`s nomenclature [15]. In this zone the carbon is constituted of rollings made of about 10 carbon planes with some porosities and some amorphous carbon, which correspond to the texture class I [15]. This interphase structure is, at room temperature, in favor of phenomenon of debonding and sliding close to the interphase/matrix interfaces, due to the nature of the bounds of carbon planes of van der Waals type; such phenomenon is also operating during creep at high temperature. It is this geometry of microstructure, as in the case of Cf – SiC, which permits both a mode I ! mode II crack deviation (Fig. 3a) and a bridging of some cracks by nano-ribbons of carbon planes (Fig. 3b), which are not an artifact as their diffraction patterns correspond to carbon: both features are nano-mechanisms which consume energy and, then, permit a certain damage tolerance.
4. Discussion and conclusion In this short paper one has shown that microscopic observations and analysis of thermo-mechanical or crept CMCs are the only way to access the different micromechanisms occurring, for example when creep tests are performed. Creep tests give only a value of some mechanical parameters and concern only a macroscopic approach. For these materials, the pyrocarbon interphases play a key role for the development of the matrix
microcracks and, consequently, permits a correct damage tolerance. To propose a creep mechanism one must accurately analyze the carbon plane orientations, their evolution, and nano- and micro-mechanisms occurring at the different microstructural scales. Materials based on selfsealing matrix exhibit a significant improvement in comparison with previous Cf – SiC or SiCf – SiC: they can be considered as a class of material for gas turbine jet engines with, for example, lifetime higher than 100 h under 170 MPa at 1473 K [17].
Acknowledgements This work has been supported by Snecma Propulsion Solide (Saint Me´dard en Jalles, France), by CNRS and Re´gion de Basse-Normandie (SD). We wish to warmly thank Drs M. Bourgeon, E. Pestourie, J.M. Rouge`s, for fruitful discussions and for providing the specimens.
References [1] Christin F. In: Krenkel W, Naslain R, Schneider H, editors. High temperature ceramic matrix composites, HT-CMC4. Weinheim, Germany: Wiley–VCH; 2001. p. 731–43. [2] Luthra KL, Corman GS. In: Krenkel W, Naslain R, Schneider H, editors. High temperature ceramic matrix composites, HT-CMC4. Weinheim, Germany: Wiley–VCH; 2001. p. 744–53. [3] Renz R, Krenkel W. Composites: from fundamentals to exploitation. ECCM 9, Brighton, UK. ECCM 9 CD ROM C 2000. IOM Communication Ltd; 4– 7 June 2000. [4] Renz R, Heidenreich B, Krenkel W, Scho¨ppach A, Richter F. In: Krenkel W, Naslain R, Schneider H, editors. High temperature ceramic matrix composites, HT-CMC4. Weinheim, Germany: Wiley–VCH; 2001. p. 839 –45. [5] Kerans RJ, Hay RS, Pagano NJ, Parthasarathy TA. Am Ceram Soc Bull 1989;68:429–42. [6] Kuntz M, Grathwohl G. Advd Engng Mater 2001;3:371– 9.
1470
G. Boitier et al. / Composites: Part A 33 (2002) 1467–1470
[7] Boitier G, Vicens J, Chermant JL. Mater Sci Engng 2000;A279: 73–80. [8] Boitier G, Vicens J, Chermant JL. Mater Sci Engng 2000;A289: 265–75. [9] Boitier G, Vicens J, Chermant JL. Mater Sci Engng 2001;A313: 53–63. [10] Boitier G. The`se de Doctorat of the University of Caen; 1997. [11] Darzens S. The`se de Doctorat of the University of Caen; 2000. [12] Chermant JL. Sil Ind 1995;60:261 –73.
[13] Kachanov L. Izv Akad Nauk SSR 1985;8:26–31. [14] Dupel P, Bourrat X, Pailler R. Carbon 1995;9:1193 –204. [15] Despre`s JF. The`se de Doctorat of the University of Pau and of Pays de l’Adour; 1993. [16] Boitier G, Vicens J, Chermant JL. Script Mater 1998;38:937–43. [17] Lamouroux F, Bouillon E, Cavalier JC, Spriet P, Habarou G. In: Krenkel W, Naslain R, Schneider H, editors. High temperature ceramic matrix composites, HT-CMC4. Weinheim, Germany: Wiley–VCH; 2001. p. 783 –8.
Microstructural investigation of interfaces in CMCs G. Boitier*, S. Darzens, J.-L. Chermant, J. Vicens LERMAT, URA CNRS 1317, ISMRA, 6 Bd Mare´chal Juin, 14050 Caen Cedex, France
Abstract This paper is focused on the importance of pyrocarbon interfaces in two types of ceramic matrix composites (Cf – SiC and SiCf – SiBC), during creep tests under argon. The development of micromechanisms which consume energy and then allow a damage tolerance, depends on the morphology of the fiber/matrix interphase, which has been investigated by TEM and HRTEM. q 2002 Elsevier Science Ltd. All rights reserved. Keywords: A. Ceramic-matrix composites (CMCs); B. Interface/interphase; TEM; B. Creep
1. Introduction Ceramic matrix composites (CMCs) are a class of materials developed for aeronautics and space applications, in a domain where superalloys cannot be used anymore. They have potential applications in structures (air intakes, structural panels with stiffners, high dimensional stability structures for mirror or antenna, etc.) or in turbines (rear frame liners, mixer flow, petals, exhaust cones, etc.), or for brakes [1 – 4]. But their mechanical behavior depends mainly on the fiber/matrix interfaces (or interphases!). They can be characterized at the macroscopic level by debonding energy, G; and frictional interface shear stress, t; but understanding the different behaviors and obtained values requires investigations at the micro- and nanoscopic scale using TEM and HRTEM. The micro- and nanostructures can be correlated to the characteristics of the G and t values, and then to the life time parameters [5 –9]. The scope of this paper is to present the complexity, the role, the evolution and the importance of interfaces and interphases in two CMCs reinforced with continuous carbon or silicon carbide fibers: Cf –SiC and SiCf – SiBC, before and after creep tests.
2. Materials and techniques Cf –SiC and SiCf – SiBC were fabricated by Snecma Propulsion Solide (St Me´dard en Jalles, France) from chemical vapor infiltration (CVI) processes in woven 2.5 D fiber architectures (ex-PAN carbon fibers or Nicalon NLM 202 silicon carbide fibers). SiBC matrix is a self-healing * Corresponding author.
multilayered matrix, based on Si – C, B – C and Si – B –C phases. All the fiber architectures have received a thin pyrolytic carbon deposit before the infiltration of the matrix. Specimens were creep tested in tension, under a partial pressure of argon (500 mbar), in a temperature domain up to 1673 K and a stress domain between 110 and 250 MPa [10, 11]. Due to the mismatch of the coefficients of thermal expansion between fibers and matrix, Cf –SiC in the asreceived state has some matrix microcracks, which are not present in SiCf –SiBC composites.
3. Results During creep these composites exhibit a typical damagecreep [10 – 12], with matrix microcracking, fiber/matrix and yarn/matrix debonding, fiber and yarn bridging, fiber and yarn pull-out and rupture. So the deformation of these CMCs can be analyzed in terms of damage mechanics, as proposed by Kachanov [13]. The microcrack network will depend on the interaction with the pyrolitic carbon (PyC) deposited on the fibers, for Cf – SiC and SiCf –SiBC, and also with the matrix multilayers for SiCf –SiBC: the role of these interphases appears as predominant. The observation of the as-received composites reveals in both cases turbostratic carbon regions, globally parallel to the fibers and the matrix (Fig. 1), as already observed [14, 15] and classified by Despre`s [15] as ‘isotropic globally oriented’ in his classification of possible PyC interphases in Cf – SiC composites. In the case of Cf – SiC the observations and analysis of the PyC interphases reveal two types of interfacial sliding mechanisms [8 –10]. At 1473 K there is the interphase rupture by debonding between two carbon
1359-835X/02/$ - see front matter q 2002 Elsevier Science Ltd. All rights reserved. PII: S 1 3 5 9 - 8 3 5 X ( 0 2 ) 0 0 1 4 7 - 1
1468
G. Boitier et al. / Composites: Part A 33 (2002) 1467–1470
Fig. 1. HRTEM micrographs of the matrix/pyrocarbon interface in a as-received Cf – SiC (a) and of the matrix/pyrocarbon/fiber interfaces in a as-received SiCf –SiBC (b).
leaves: then the interfacial sliding can be assimilated to a dry friction between two rough solids. Close to the fibers and the matrix the structure of the PyC is turbostratic type, with atomic carbon planes parallel to the fiber direction. That has also clearly been evidenced by SEM observations of the surface of the fibers: that surface is very rough (see the small insert in Fig. 2a). Although at 1473 K the fiber/matrix interface is radially in compression, under stress some lenticular pores appear (Fig. 2a). These lenticular pores due to local decohesions between carbon planes are, in fact, the nuclei for the microcrack development. When the debonding is not total, carbon ribbons bridge the two parts of the microcracks. At 1673 K, the PyC interphase is a little degraded over about 100 nm from the matrix; there are a disappearing of the previous anisotropic texture, and a certain amorphization; then the interfacial sliding can be assimilated to a viscous flow (Fig. 2b). That was also confirmed by SEM observations: the fiber surfaces are smooth [8 –10] (see the small insert in Fig. 2b). If observations are performed at a higher scale one notes that the nanotexture of the carbon fibers increases with the test temperature. Consequently it leads to an increase of the local molecular orientation of oriented volumes of carbon planes parallel to the fibers. Although this phenomenon has been called ‘carbon fiber nanocreep’ [16], this contribution of the fibers and matrix to the macroscopic deformation is negligible, but it can be considered as the nuclei of the damage process. In the case of SiCf – SiBC same type of features are observed, but a little more complex due to the existence of the different matrix layers based on the Si –B – C system. In addition to the damage observed in Cf – SiC composites, there are also matrix microcrack deviations via some matrix multilayers where carbon exists [11]: this carbon is located at the interfaces of specific matrix layers and presents a thin
Fig. 2. HRTEM and SEM micrographs of the pyrocarbon interphase in Cf –SiC creep tested specimen under 220 MPa and in Ar: (a) with the presence of some lenticular pores at 1473 K; (b) and the presence of carbon rollings at 1673 K. Inserts correspond to SEM images of the surface features of the SiCf fibers at these two temperatures.
G. Boitier et al. / Composites: Part A 33 (2002) 1467–1470
1469
Fig. 3. TEM micrographs of the pyrocarbon interphase showing a microcrack deviation mode I ! mode II in a longitudinal section (parallel to the stress direction) (a), and a bridging of a microcrack by carbon ribbons (b), for SiCf –SiBC specimens creep tested at 1473 K, under 200 MPa and in Ar.
layer of PyC, from 10 to 100 nm thick, and oriented globally parallel to these interfaces. The larger the interphase is, the more pronounced the debonding between the matrix layers appears. In the case of the pyrocarbon layer (between the silicon carbide fibers and the matrix), close to the matrix the carbon structure is turbostratic type over a thickness of about 12 – 15 nm, i.e. 30– 35 carbon atomic planes. Further away, the carbon structure is globally isotropic oriented, according to the Despre`s nomenclature [15]. In this zone the carbon is constituted of rollings made of about 10 carbon planes with some porosities and some amorphous carbon, which correspond to the texture class I [15]. This interphase structure is, at room temperature, in favor of phenomenon of debonding and sliding close to the interphase/matrix interfaces, due to the nature of the bounds of carbon planes of van der Waals type; such phenomenon is also operating during creep at high temperature. It is this geometry of microstructure, as in the case of Cf – SiC, which permits both a mode I ! mode II crack deviation (Fig. 3a) and a bridging of some cracks by nano-ribbons of carbon planes (Fig. 3b), which are not an artifact as their diffraction patterns correspond to carbon: both features are nano-mechanisms which consume energy and, then, permit a certain damage tolerance.
4. Discussion and conclusion In this short paper one has shown that microscopic observations and analysis of thermo-mechanical or crept CMCs are the only way to access the different micromechanisms occurring, for example when creep tests are performed. Creep tests give only a value of some mechanical parameters and concern only a macroscopic approach. For these materials, the pyrocarbon interphases play a key role for the development of the matrix
microcracks and, consequently, permits a correct damage tolerance. To propose a creep mechanism one must accurately analyze the carbon plane orientations, their evolution, and nano- and micro-mechanisms occurring at the different microstructural scales. Materials based on selfsealing matrix exhibit a significant improvement in comparison with previous Cf – SiC or SiCf – SiC: they can be considered as a class of material for gas turbine jet engines with, for example, lifetime higher than 100 h under 170 MPa at 1473 K [17].
Acknowledgements This work has been supported by Snecma Propulsion Solide (Saint Me´dard en Jalles, France), by CNRS and Re´gion de Basse-Normandie (SD). We wish to warmly thank Drs M. Bourgeon, E. Pestourie, J.M. Rouge`s, for fruitful discussions and for providing the specimens.
References [1] Christin F. In: Krenkel W, Naslain R, Schneider H, editors. High temperature ceramic matrix composites, HT-CMC4. Weinheim, Germany: Wiley–VCH; 2001. p. 731–43. [2] Luthra KL, Corman GS. In: Krenkel W, Naslain R, Schneider H, editors. High temperature ceramic matrix composites, HT-CMC4. Weinheim, Germany: Wiley–VCH; 2001. p. 744–53. [3] Renz R, Krenkel W. Composites: from fundamentals to exploitation. ECCM 9, Brighton, UK. ECCM 9 CD ROM C 2000. IOM Communication Ltd; 4– 7 June 2000. [4] Renz R, Heidenreich B, Krenkel W, Scho¨ppach A, Richter F. In: Krenkel W, Naslain R, Schneider H, editors. High temperature ceramic matrix composites, HT-CMC4. Weinheim, Germany: Wiley–VCH; 2001. p. 839 –45. [5] Kerans RJ, Hay RS, Pagano NJ, Parthasarathy TA. Am Ceram Soc Bull 1989;68:429–42. [6] Kuntz M, Grathwohl G. Advd Engng Mater 2001;3:371– 9.
1470
G. Boitier et al. / Composites: Part A 33 (2002) 1467–1470
[7] Boitier G, Vicens J, Chermant JL. Mater Sci Engng 2000;A279: 73–80. [8] Boitier G, Vicens J, Chermant JL. Mater Sci Engng 2000;A289: 265–75. [9] Boitier G, Vicens J, Chermant JL. Mater Sci Engng 2001;A313: 53–63. [10] Boitier G. The`se de Doctorat of the University of Caen; 1997. [11] Darzens S. The`se de Doctorat of the University of Caen; 2000. [12] Chermant JL. Sil Ind 1995;60:261 –73.
[13] Kachanov L. Izv Akad Nauk SSR 1985;8:26–31. [14] Dupel P, Bourrat X, Pailler R. Carbon 1995;9:1193 –204. [15] Despre`s JF. The`se de Doctorat of the University of Pau and of Pays de l’Adour; 1993. [16] Boitier G, Vicens J, Chermant JL. Script Mater 1998;38:937–43. [17] Lamouroux F, Bouillon E, Cavalier JC, Spriet P, Habarou G. In: Krenkel W, Naslain R, Schneider H, editors. High temperature ceramic matrix composites, HT-CMC4. Weinheim, Germany: Wiley–VCH; 2001. p. 783 –8.