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In situ characterization of tensile damage behavior of a plain-woven fiber-reinforced polymer-derived ceramic composite
Materials Letters 57 (2003) 1571 – 1578 www.elsevier.com/locate/matlet
In situ characterization of tensile damage behavior of a plain-woven fiber-reinforced polymer-derived ceramic composite Yufu Liu *, Yoshihisa Tanaka National Institute for Materials Science 1-2-1 Sengen, Tsukuba 305-0047, Japan Received 25 March 2002; accepted 15 July 2002
Abstract A tensile test of a plain-woven SiC composite, manufactured with the polymer infiltration and pyrolysis method, was carried out. Various damage events and their sequences were characterized by in situ SEM observations. D 2002 Elsevier Science B.V. All rights reserved. Keywords: In situ characterization; Tensile damage behavior; Ceramic composite
1. Introduction The failure mechanisms of fiber-reinforced ceramic matrix composites (CMC) are known to involve complicated damage modes and are affected by many factors such as the reinforcement architecture, processing routes (liquid or gas) and conditions, and fiber – matrix interphases. If the interface between the fiber and matrix in a unidirectionally aligned continuous fiberreinforced system is weakly bonded, interface debonding and subsequent sliding as a major source of energy dissipation occur and play an important role in the fracture process ([1] and references therein). In two-dimensional woven-fiber composites, transverse fiber bundles are weak zones when the loading is applied along the longitudinal fiber direction; processing-induced inter- and intra-lamina flaws and their distribution affect matrix cracking initiation behavior *
Corresponding author. Tel.: +81-298-59-2447; fax: +81-29859-2401. E-mail address: [email protected] (Y. Liu).
[2]. It was reported [3] that a two-dimensional woven Nicalon fiber-reinforced SiCON matrix composite can provide high damage-tolerance because crossover points of longitudinal and transverse fiber bundles may be debonded and then experience sliding possibly due to strain mismatches between the bundles, even when the fiber – matrix interface is strongly bonded. The damage modes are further complicated in the three-dimensional reinforcement system [2]. Typical stress – strain curves of these composites usually show a linear response initially, followed by a long nonlinear range and finally terminated with reinforcement failure. CMCs manufactured by chemical vapor infiltration (CVI) have been investigated extensively and the damage mechanisms are well documented in the literature. Meanwhile, CMCs may also be fabricated by the liquid route using polymer infiltration and pyrolysis (PIP). CVI is powerful for obtaining highly densified materials with spatial homogeneity [4], while PIP usually yields less dense composites. For example, a well-controlled CVI CMC may contain
0167-577X/02/$ - see front matter D 2002 Elsevier Science B.V. All rights reserved. doi:10.1016/S0167-577X(02)01034-0
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Y. Liu, Y. Tanaka / Materials Letters 57 (2003) 1571–1578
less than 5% porosity, but a PIP material generally has about 8– 23% of porosity depending on the number of PIP infiltration cycles [5]. In the previous studies, composites based on PIP have received less attention despite its relatively low cost. In addition, the interlamina damage in these studies is mixed to various other damage patterns and this makes it difficult to observe real damage evolution behavior. In this study, tensile testing of a cross-woven one-ply SiC fiberreinforced SiC composite fabricated by PIP is carried out in the scanning electron microscope (SEM) in order to observe the cracking evolution on the side and upper surfaces of the composite. The use of one-ply cross-woven fibers is to exclude inter-lamina flaws and thus provides a means to separate damage contribution. The tensile fracture behavior of the composite associated with various damage events and their sequences is characterized by in situ observation and measurement.
Fig. 2. Typical stress – strain and unloading – reloading hysterisis curves obtained in the experiment. Also shown are the various damage events and their sequences.
2. Experimental procedure BN-coated-HiNicalon SiC fiber ( Nippon Carbon, Tokyo, Japan) was used as the reinforcement. During
the PIP process, one ply of the plain-woven fibers was infiltrated with an organic polymer. The polymer was heated to fairly high temperatures of about 900 jC
Fig. 1. Optical micrograph showing the microstructure of the PIP composite.
Y. Liu, Y. Tanaka / Materials Letters 57 (2003) 1571–1578
and several infiltrations were used to densify the composites [6]. Fig. 1 shows an optical micrograph of the as-received material consisting of two-dimensional woven-fiber bundles, with each bundle containing 500 fibers. Processing-induced voids give an average void volume fraction of about 16%; the nominal fiber volume fraction is approximately 43% and the remainder is the matrix [6]. Smooth rectangular type specimens 50 mm long, 10 mm wide, and 0.42 mm thick were tested. Aluminum tabs were glued on the two ends of the specimen
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and the specimen was mounted on a motor-driven testing jig with a load cell in SEM [7]. The specimen was loaded along the longitudinal fiber bundle direction. The tensile strain was measured by strain gauges 3 mm long and 2 mm wide. Two surfaces (upper and side) of the specimen were polished for observation with diamond paste up to 0.25 Am finish. Tensile loading was applied step by step and the tests were frequently interrupted for direct in situ observation. Unloading and reloading tests were also conducted to record damage accumulation.
Fig. 3. SEM micrographs showing fracture events at applied stresses of 27.5 MPa (a), 52.5 MPa (b), 116.2 MPa (c) and 116.2 MPa (d). Arrows in (a) and (b) indicate fracture initiation.
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Y. Liu, Y. Tanaka / Materials Letters 57 (2003) 1571–1578
3. Results and discussion Fig. 2 shows typical stress –strain response curves and damage patterns identified in the experiment. This curve has a short linear range at the beginning and is essentially nonlinear later. Permanent strains were observed upon complete unloading; the permanent strain increased with the increase of the unloading stress level and the unloading –reloading hysteresis loop also became wider. The various damage levels and modes presented in Fig. 2 will be elaborated and discussed in the following sections. During the experiment, two types of the initial cracking event were observed on the upper surface: one was originated from processing-induced voids (Fig. 3(a)), while the other one is fiber – matrix interface debonding in transverse fiber bundles (Fig. 3(b)). Which one occurred earlier was found to be rather specimen-dependent. Upon further loading, the number of cracks initiated from processing-induced voids and transverse bundles increased (Fig. 3(c) and (d)). On the side surface, the crack initiated from the processing void propagated (Fig. 4(a)) and partial debonding of the transverse fiber bundle occurred (Fig. 4(b)) when the loading was further increased. Following all these damages were fiber breaks in longitudinal fiber bundles and further extension of the matrix cracking from the transverse fiber bundle which was then retarded by the longitudinal fiber bundle (Fig. 5). The fiber breaks occurred in a cumulative manner because of the scattered strength of the fibers and the stress concentration in the bending longitudinal fibers. It should be pointed out that almost no debonding crack at the fiber –matrix interface in the longitudinal fiber bundle was observed. Final catastrophic failure of the composite was caused by fiber breaks in longitudinal fiber bundles (Fig. 6). The fracture surface is brush-like with a small amount of fiber pullout, indicating again that there is strong bonding between the fiber and matrix in the longitudinal fiber bundle. The crack initiation behavior in this woven-fiber composite showed no much difference from that reported previously [2]. However, the fracture process was different from the well-documented mechanisms of matrix cracking followed by the fiber –matrix debonding and deflection and finally fiber breaks and pullout.
Fig. 4. Matrix cracking and partial debonding at the longitudinal and transverse fiber bundle interface. The loading stress for (a) was 92.3 MPa and for (b) 124.1 MPa.
The stress –strain hysteresis showed some features typical of ductile materials. From the above observations, the nonlinear stress – strain response is believed to result from the matrix cracking initiated the processing voids and matrix – fiber interface debonding in the transverse bundle and progressive debonding between longitudinal and transverse fiber bundle. In order to prove this, further detailed observation of the damage modes at various locations of the upper and side surfaces was attempted. Fig. 7(a) and (b) shows the matrix crack opening at a loading stress of 92.3 MPa and closure behavior
Y. Liu, Y. Tanaka / Materials Letters 57 (2003) 1571–1578
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Fig. 5. Fiber breakage in the longitudinal fiber bundle and retardation of a transverse crack by the longitudinal fiber bundle. The loading stress was 124.1 MPa.
on complete unloading, respectively. In Fig. 7(b), partial crack opening and closure were observed; the partial crack opening contributed to the perma-
Fig. 6. Fracture surface showing fiber pullout in the longitudinal fiber bundle. The final failure stress was about 156 MPa.
nent deformation, while the crack closure caused asperities and wear during unloading and reloading. A higher magnification micrograph shown in Fig. 7(c) indicates the crack opening at complete unloading is about. Fig. 8(a) and (b) present typical sidesurface damage patterns showing the crack opening at a loading stress of 124.1 MPa and closure behavior on complete unloading, respectively. The rough crack path in the transverse bundle contributes to nonlinear unloading – reloading hysteresis and is believed to be a source of permanent ‘plastic’ strain accumulation. Past studies have not paid due attention to this type of asperities and wear. The crack closure occurring upon complete unloading (Figs. 7 and 8) is noteworthy. In comparison with the past reports ([1,8] and references therein) where the ductility and toughness of the composite are attributed to the fiber – matrix interface debonding, sliding, wear and change of friction in the longitudinal fiber bundles, the present study demonstrates that even there is no significant debonding crack at the fiber – matrix interface in the longitudinal fiber bundle, ‘ductility’ and high toughness of the com-
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Y. Liu, Y. Tanaka / Materials Letters 57 (2003) 1571–1578
Fig. 7. Transverse crack opening behavior at an applied stress of 92.3 MPa (a), at complete unloading (b), and higher magnification of the dotted line rectangle (c).
posite can still be obtained by matrix cracking from the processing voids and matrix – fiber interface debonding in the transverse bundle and progressive debonding between longitudinal and transverse bundles. These may be regarded as a feature of the present PIP CMC.
4. Conclusions (1) A novel tensile test of a cross-woven one-ply SiC fiber-reinforced SiC composite manufactured with the polymer infiltration and pyrolysis (PIP) method was carried out in a scanning electron micro-
Y. Liu, Y. Tanaka / Materials Letters 57 (2003) 1571–1578
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Fig. 8. Matrix crack and partial debonding crack opening at the longitudinal and transverse fiber bundle interface. Micrographs (a) and (b) were taken at an applied stress of 124.1 MPa and at complete unloading, respectively; (a) and (b) roughly correspond to the dotted line rectangle of (c).
scope (SEM). The various damage events and their sequences were characterized by in situ SEM observation. (2) Matrix cracking from the processing voids and matrix – fiber interface debonding in the transverse bundle and progressive debonding between longitudinal and transverse fiber bundle play an important role in the composite ductility and toughness enhancement
even when the fiber – matrix interface in the longitudinal fiber bundle is strongly bonded.
Acknowledgements The authors would like to thank Dr. C. Masuda of National Institute for Materials Science and Prof. Y.
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Y. Liu, Y. Tanaka / Materials Letters 57 (2003) 1571–1578
Kagawa of The University of Tokyo for the useful comments in the course of the present work.
References [1] Y.F. Liu, Advances in Condensed Matter and Materials Research, Nova Science Publishers, New York, USA, 2001, pp. 71 – 94. [2] T.-W. Chou, Microstructural design of fiber composites, Cambridge Univ. Press, Great Britain, 1992, pp. 285 – 440.
[3] N. Chawla, Y.K. Tur, J.W. Holmes, J.R. Barber, J. Am. Ceram. Soc. 81 (1998) 1221 – 1230. [4] N. Yoshikawa, J.W. Evans, J. Am. Ceram. Soc. 85 (2002) 1485 – 1491. [5] M. Takeda, Y. Kagawa, S. Matsuno, Y. Imai, H. Ichikawa, J. Am. Ceram. Soc. 82 (1999) 1579 – 1581. [6] Private communication with Nippon Carbon, Tokyo, Japan. [7] Y.-F. Liu, Y. Tanaka, C. Masuda, Acta Mater. 46 (1998) 5237 – 5247. [8] B. Budiansky, A.G. Evans, J.W. Hutchinson, Int. J. Solids Struct. 32 (1995) 315 – 328.
In situ characterization of tensile damage behavior of a plain-woven fiber-reinforced polymer-derived ceramic composite Yufu Liu *, Yoshihisa Tanaka National Institute for Materials Science 1-2-1 Sengen, Tsukuba 305-0047, Japan Received 25 March 2002; accepted 15 July 2002
Abstract A tensile test of a plain-woven SiC composite, manufactured with the polymer infiltration and pyrolysis method, was carried out. Various damage events and their sequences were characterized by in situ SEM observations. D 2002 Elsevier Science B.V. All rights reserved. Keywords: In situ characterization; Tensile damage behavior; Ceramic composite
1. Introduction The failure mechanisms of fiber-reinforced ceramic matrix composites (CMC) are known to involve complicated damage modes and are affected by many factors such as the reinforcement architecture, processing routes (liquid or gas) and conditions, and fiber – matrix interphases. If the interface between the fiber and matrix in a unidirectionally aligned continuous fiberreinforced system is weakly bonded, interface debonding and subsequent sliding as a major source of energy dissipation occur and play an important role in the fracture process ([1] and references therein). In two-dimensional woven-fiber composites, transverse fiber bundles are weak zones when the loading is applied along the longitudinal fiber direction; processing-induced inter- and intra-lamina flaws and their distribution affect matrix cracking initiation behavior *
Corresponding author. Tel.: +81-298-59-2447; fax: +81-29859-2401. E-mail address: [email protected] (Y. Liu).
[2]. It was reported [3] that a two-dimensional woven Nicalon fiber-reinforced SiCON matrix composite can provide high damage-tolerance because crossover points of longitudinal and transverse fiber bundles may be debonded and then experience sliding possibly due to strain mismatches between the bundles, even when the fiber – matrix interface is strongly bonded. The damage modes are further complicated in the three-dimensional reinforcement system [2]. Typical stress – strain curves of these composites usually show a linear response initially, followed by a long nonlinear range and finally terminated with reinforcement failure. CMCs manufactured by chemical vapor infiltration (CVI) have been investigated extensively and the damage mechanisms are well documented in the literature. Meanwhile, CMCs may also be fabricated by the liquid route using polymer infiltration and pyrolysis (PIP). CVI is powerful for obtaining highly densified materials with spatial homogeneity [4], while PIP usually yields less dense composites. For example, a well-controlled CVI CMC may contain
0167-577X/02/$ - see front matter D 2002 Elsevier Science B.V. All rights reserved. doi:10.1016/S0167-577X(02)01034-0
1572
Y. Liu, Y. Tanaka / Materials Letters 57 (2003) 1571–1578
less than 5% porosity, but a PIP material generally has about 8– 23% of porosity depending on the number of PIP infiltration cycles [5]. In the previous studies, composites based on PIP have received less attention despite its relatively low cost. In addition, the interlamina damage in these studies is mixed to various other damage patterns and this makes it difficult to observe real damage evolution behavior. In this study, tensile testing of a cross-woven one-ply SiC fiberreinforced SiC composite fabricated by PIP is carried out in the scanning electron microscope (SEM) in order to observe the cracking evolution on the side and upper surfaces of the composite. The use of one-ply cross-woven fibers is to exclude inter-lamina flaws and thus provides a means to separate damage contribution. The tensile fracture behavior of the composite associated with various damage events and their sequences is characterized by in situ observation and measurement.
Fig. 2. Typical stress – strain and unloading – reloading hysterisis curves obtained in the experiment. Also shown are the various damage events and their sequences.
2. Experimental procedure BN-coated-HiNicalon SiC fiber ( Nippon Carbon, Tokyo, Japan) was used as the reinforcement. During
the PIP process, one ply of the plain-woven fibers was infiltrated with an organic polymer. The polymer was heated to fairly high temperatures of about 900 jC
Fig. 1. Optical micrograph showing the microstructure of the PIP composite.
Y. Liu, Y. Tanaka / Materials Letters 57 (2003) 1571–1578
and several infiltrations were used to densify the composites [6]. Fig. 1 shows an optical micrograph of the as-received material consisting of two-dimensional woven-fiber bundles, with each bundle containing 500 fibers. Processing-induced voids give an average void volume fraction of about 16%; the nominal fiber volume fraction is approximately 43% and the remainder is the matrix [6]. Smooth rectangular type specimens 50 mm long, 10 mm wide, and 0.42 mm thick were tested. Aluminum tabs were glued on the two ends of the specimen
1573
and the specimen was mounted on a motor-driven testing jig with a load cell in SEM [7]. The specimen was loaded along the longitudinal fiber bundle direction. The tensile strain was measured by strain gauges 3 mm long and 2 mm wide. Two surfaces (upper and side) of the specimen were polished for observation with diamond paste up to 0.25 Am finish. Tensile loading was applied step by step and the tests were frequently interrupted for direct in situ observation. Unloading and reloading tests were also conducted to record damage accumulation.
Fig. 3. SEM micrographs showing fracture events at applied stresses of 27.5 MPa (a), 52.5 MPa (b), 116.2 MPa (c) and 116.2 MPa (d). Arrows in (a) and (b) indicate fracture initiation.
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Y. Liu, Y. Tanaka / Materials Letters 57 (2003) 1571–1578
3. Results and discussion Fig. 2 shows typical stress –strain response curves and damage patterns identified in the experiment. This curve has a short linear range at the beginning and is essentially nonlinear later. Permanent strains were observed upon complete unloading; the permanent strain increased with the increase of the unloading stress level and the unloading –reloading hysteresis loop also became wider. The various damage levels and modes presented in Fig. 2 will be elaborated and discussed in the following sections. During the experiment, two types of the initial cracking event were observed on the upper surface: one was originated from processing-induced voids (Fig. 3(a)), while the other one is fiber – matrix interface debonding in transverse fiber bundles (Fig. 3(b)). Which one occurred earlier was found to be rather specimen-dependent. Upon further loading, the number of cracks initiated from processing-induced voids and transverse bundles increased (Fig. 3(c) and (d)). On the side surface, the crack initiated from the processing void propagated (Fig. 4(a)) and partial debonding of the transverse fiber bundle occurred (Fig. 4(b)) when the loading was further increased. Following all these damages were fiber breaks in longitudinal fiber bundles and further extension of the matrix cracking from the transverse fiber bundle which was then retarded by the longitudinal fiber bundle (Fig. 5). The fiber breaks occurred in a cumulative manner because of the scattered strength of the fibers and the stress concentration in the bending longitudinal fibers. It should be pointed out that almost no debonding crack at the fiber –matrix interface in the longitudinal fiber bundle was observed. Final catastrophic failure of the composite was caused by fiber breaks in longitudinal fiber bundles (Fig. 6). The fracture surface is brush-like with a small amount of fiber pullout, indicating again that there is strong bonding between the fiber and matrix in the longitudinal fiber bundle. The crack initiation behavior in this woven-fiber composite showed no much difference from that reported previously [2]. However, the fracture process was different from the well-documented mechanisms of matrix cracking followed by the fiber –matrix debonding and deflection and finally fiber breaks and pullout.
Fig. 4. Matrix cracking and partial debonding at the longitudinal and transverse fiber bundle interface. The loading stress for (a) was 92.3 MPa and for (b) 124.1 MPa.
The stress –strain hysteresis showed some features typical of ductile materials. From the above observations, the nonlinear stress – strain response is believed to result from the matrix cracking initiated the processing voids and matrix – fiber interface debonding in the transverse bundle and progressive debonding between longitudinal and transverse fiber bundle. In order to prove this, further detailed observation of the damage modes at various locations of the upper and side surfaces was attempted. Fig. 7(a) and (b) shows the matrix crack opening at a loading stress of 92.3 MPa and closure behavior
Y. Liu, Y. Tanaka / Materials Letters 57 (2003) 1571–1578
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Fig. 5. Fiber breakage in the longitudinal fiber bundle and retardation of a transverse crack by the longitudinal fiber bundle. The loading stress was 124.1 MPa.
on complete unloading, respectively. In Fig. 7(b), partial crack opening and closure were observed; the partial crack opening contributed to the perma-
Fig. 6. Fracture surface showing fiber pullout in the longitudinal fiber bundle. The final failure stress was about 156 MPa.
nent deformation, while the crack closure caused asperities and wear during unloading and reloading. A higher magnification micrograph shown in Fig. 7(c) indicates the crack opening at complete unloading is about. Fig. 8(a) and (b) present typical sidesurface damage patterns showing the crack opening at a loading stress of 124.1 MPa and closure behavior on complete unloading, respectively. The rough crack path in the transverse bundle contributes to nonlinear unloading – reloading hysteresis and is believed to be a source of permanent ‘plastic’ strain accumulation. Past studies have not paid due attention to this type of asperities and wear. The crack closure occurring upon complete unloading (Figs. 7 and 8) is noteworthy. In comparison with the past reports ([1,8] and references therein) where the ductility and toughness of the composite are attributed to the fiber – matrix interface debonding, sliding, wear and change of friction in the longitudinal fiber bundles, the present study demonstrates that even there is no significant debonding crack at the fiber – matrix interface in the longitudinal fiber bundle, ‘ductility’ and high toughness of the com-
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Y. Liu, Y. Tanaka / Materials Letters 57 (2003) 1571–1578
Fig. 7. Transverse crack opening behavior at an applied stress of 92.3 MPa (a), at complete unloading (b), and higher magnification of the dotted line rectangle (c).
posite can still be obtained by matrix cracking from the processing voids and matrix – fiber interface debonding in the transverse bundle and progressive debonding between longitudinal and transverse bundles. These may be regarded as a feature of the present PIP CMC.
4. Conclusions (1) A novel tensile test of a cross-woven one-ply SiC fiber-reinforced SiC composite manufactured with the polymer infiltration and pyrolysis (PIP) method was carried out in a scanning electron micro-
Y. Liu, Y. Tanaka / Materials Letters 57 (2003) 1571–1578
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Fig. 8. Matrix crack and partial debonding crack opening at the longitudinal and transverse fiber bundle interface. Micrographs (a) and (b) were taken at an applied stress of 124.1 MPa and at complete unloading, respectively; (a) and (b) roughly correspond to the dotted line rectangle of (c).
scope (SEM). The various damage events and their sequences were characterized by in situ SEM observation. (2) Matrix cracking from the processing voids and matrix – fiber interface debonding in the transverse bundle and progressive debonding between longitudinal and transverse fiber bundle play an important role in the composite ductility and toughness enhancement
even when the fiber – matrix interface in the longitudinal fiber bundle is strongly bonded.
Acknowledgements The authors would like to thank Dr. C. Masuda of National Institute for Materials Science and Prof. Y.
1578
Y. Liu, Y. Tanaka / Materials Letters 57 (2003) 1571–1578
Kagawa of The University of Tokyo for the useful comments in the course of the present work.
References [1] Y.F. Liu, Advances in Condensed Matter and Materials Research, Nova Science Publishers, New York, USA, 2001, pp. 71 – 94. [2] T.-W. Chou, Microstructural design of fiber composites, Cambridge Univ. Press, Great Britain, 1992, pp. 285 – 440.
[3] N. Chawla, Y.K. Tur, J.W. Holmes, J.R. Barber, J. Am. Ceram. Soc. 81 (1998) 1221 – 1230. [4] N. Yoshikawa, J.W. Evans, J. Am. Ceram. Soc. 85 (2002) 1485 – 1491. [5] M. Takeda, Y. Kagawa, S. Matsuno, Y. Imai, H. Ichikawa, J. Am. Ceram. Soc. 82 (1999) 1579 – 1581. [6] Private communication with Nippon Carbon, Tokyo, Japan. [7] Y.-F. Liu, Y. Tanaka, C. Masuda, Acta Mater. 46 (1998) 5237 – 5247. [8] B. Budiansky, A.G. Evans, J.W. Hutchinson, Int. J. Solids Struct. 32 (1995) 315 – 328.