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Weak interface dominated high temperature fracture strength of carbon fiber reinforced mullite matrix composites
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Weak interface dominated high temperature fracture strength of carbon fiber reinforced mullite matrix composites L.W. Yang b,∗ , H.T. Liu a,∗ , R. Jiang a , X. Sun a , W.G. Mao c , H.F. Cheng a , J.M. Molina-Aldareguía b a
Science and Technology on Advanced Ceramic Fibers and Composites Laboratory, National University of Defense Technology, Changsha 410073, China IMDEA Materials Institute, c/Eric Kandel 2, Getafe 28906, Madrid, Spain c School of Materials Science and Engineering, Xiangtan University, Xiangtan 411105, China b
a r t i c l e
i n f o
Article history: Received 23 October 2016 Received in revised form 14 March 2017 Accepted 17 March 2017 Available online xxx Keywords: A. Ceramic matrix composites B. Interfaces C. Fracture mechanics/toughness D. Strength
a b s t r a c t Weak fiber/matrix interface dominates the toughening properties of ceramic matrix composites. This paper reports a novel sol-gel fabricated carbon fiber reinforced mullite matrix composite, in which the fiber/matrix interface was inherently weak in shear properties (∼25 MPa), measured in-situ by fiber push-in tests. The interface microstructure was chemically sharp, characterized by transmission electron microscopy. The outcome of the weak interface was the full trigger of the toughening mechanisms like crack deflection, etc., leading to significant enhancement of the fracture toughness of the composite √ (∼12 MPa m), measured by single edged notch beam method. Finally, due to the weak fiber/matrix interface and large thermal expansion mismatch of the fiber and matrix, the high temperature fracture strength was enhanced in the temperature range from 25 to 1200 ◦ C, which is attributed to the enhancement of the interfacial property at elevated temperatures that favors better load transfers between composite constituents. © 2017 Published by Elsevier Ltd.
1. Introduction Ceramic matrix composites reinforced by continuous high strength fibers are promising materials with excellent thermal mechanical properties, e.g. high strength, improved toughness, low density, etc. [1,2]. The toughness of CMCs is still under study due to the concerns of the catastrophic brittle failures of CMCs in response to external stresses. Various methods have been developed to enhance the fracture resistance of the CMCs by tailoring the fiber/matrix interface. Among them, one common method is to introduce a compliant (ex-situ) interphase between the fiber and matrix by chemical vapor deposition, liquid phase processes etc., which are time-consuming and expensive [3,4]. Recently, some insitu methods have been proposed in carbon fiber reinforced SiC matrix (Cf/SiC) composites to fabricate the interphase directly by fiber heat treatment or by matrix fabrication [5,6]. These in-situ formed interphases have shown many advantages over the typical ex-situ interphases. Firstly, no precursors would be needed, avoiding the mechanical degradation of the carbon fibers; secondly, the
∗ Corresponding authors. E-mail addresses: [email protected] (L.W. Yang), [email protected] (H.T. Liu).
fabrication processes would be much easier and cheaper than most ex-situ processes; and finally, the in-situ interphase could be grown homogeneously, especially in the interior of three-dimensional fiber preforms, that are difficult to access by ex-situ coating methods. Despite that, an even more efficient, but challenging way to enhance the fracture resistance of CMCs is to introduce weak interface instead of compliant interphase directly during the composite fabrication process. However, the potential interfacial reactions between the fibers and the matrix during high temperature fabrications process would bring chemically controlled interfaces in many typical CMCs, thus leading to pretty strong interfacial interactions [7–10]. For example, in those SiC matrix composites that are fabricated by precursor infiltration and pyrolysis (PIP) of polycarbosilane (PCS), the Si H groups in PCS can be easily embedded by the O groups on fiber surfaces at high temperatures, forming O Si O like interphase that can bond the fiber/matrix interface tightly [9,11,12]. Therefore, proper choices of matrix and fiber materials to prevent potential interface reactions can promisingly enable the formation of weak interfaces in CMCs, without the introduction of ex-situ or in-situ interphases. As is known that mullite is a typical ceramic that can be hardly reacted with carbon materials due to the high carbo-thermal reduction reaction temperature (1800 ◦ C) [13], which could bring
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Fig. 1. (a) Surface morphology of the as-received T300 carbon fiber; (b) Polished cross-section of Cf/Mu; (c) BF-TEM of a TEM lamellae at the interface zone in Cf/Mu; (d) HR-TEM of the interface zone; (e) SAD pattern of the carbon fiber; (f) SAD pattern of the mullite matrix.
an inherently chemical sharp mullite/carbon interface. To take full advantage of this peculiarity, carbon fiber reinforced mullite matrix (Cf/Mu) composites have been developed decades ago [14]. Despite this, the interfacial properties in the Cf/Mu composites and their dependence on the interfacial microstructure and the macro fracture resistance of the Cf/Mu composites, especially at high temperatures, are still not fully explored. This information is critical not only from the fundamental viewpoint, but from the engineering perspectives since most CMCs are served in harsh environments. For the sake of filling the research gap, a novel Cf/Mu composite was fabricated in this work by sol-gel process. The fiber-matrix interface was characterized by transmission electron microscopy (TEM). The shear strength of the fiber-matrix interface was measured insitu by fiber push-in method. The mechanical fracture toughness of the composites was measured by single edged notch beam (SENB) test, and a digital image correction (DIC) method was employed to track the crack propagations to study the toughening mechanisms. Finally, the effect of the resulting interfacial property on the high temperature fracture strength of the Cf/Mu composite was studied by high temperature three-point bend (TPB) techniques in a wide temperature range from 25 to 1200 ◦ C in vaccum. 2. Experimental procedures PAN based 3 dimensional (3D) 3K-T300 carbon fiber fabrics were employed as the reinforcement. The fabric braid angle was 22◦
and the fiber volume fraction was ∼43%. The Diphasic mullite sols (AS12, from Suzhou Nanodispersions Ltd, China) prepared by comixing of silica and alumina sols were used as precursor of the mullite matrix. 3D Cf/Mu composites were fabricated by sol-gel process [15]. This method is advantageous over those traditional densification methods (slurry infiltration etc.) to achieve a denser mullite matrix at a comparably lower temperature, which is important to minimize the thermal-erosions and corrosions of the carbon fibers during matrix fabrication process [16–19]. The details of the sol-gel process are as follows: firstly, the mullite sols were infiltrated into carbon fiber preforms in a vacuum apparatus at room temperature for 4 h. After that, the infiltrated preform was sintered in Ar at 1200 ◦ C for 1 h after dried at 80 ◦ C for 12 h to gelate the mullite sols. Finally, the process was repeated for at least 14 times to obtain dense Cf/Mu composites with the relative composite weight increase less than 1% during subsequent sol-gel cycles. The morphologies of carbon fibers and Cf/Mu composites were observed by scanning electron microscope (SEM, Hitachi FEG S4800). The microstructure of the fiber/matrix interface zone was characterized by TEM (JEOL JEM 2100) with 200 kV accelerating voltage. The electron transparent TEM lamella with <100 nm in thickness was milled in a Helios 600i focused ion beam (FIB) system. The composite interface shear strength was evaluated by fiber push-in tests within a TI 950 TriboindenterTM nanoindentation system using a flat punch of 3 m in diameter. More than 8 push-in tests were carried out on selected fibers until fiber debonding. The
Please cite this article in press as: L.W. Yang, et al., Weak interface dominated high temperature fracture strength of carbon fiber reinforced mullite matrix composites, J Eur Ceram Soc (2017), http://dx.doi.org/10.1016/j.jeurceramsoc.2017.03.033
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Table 1 Review of the interfacial shear strength in the typical CMCs. (CVI: chemical vapor infiltration; PIP: precursor infiltration pyrolysis). Composites
Processing
Method
(MPa)
Ref
Cf/Mu SiCf/SiC SiCf/BN/SiC SiC/PyC/SiC SiCf/Mu Muf/ZrO2 /Al2 O3 Al2 O3 /LaPO4 /Al2 O3 Cf/SiC Cf/C/SiC Muf/SiC Muf/C/SiC
Sol-gel CVI CVI CVI Hot Press Cold Press + Sintering Sintering PIP CVI PIP PIP
Push in Push in Push out Push out Pull out Push in Push out Push in Push out Push in Push in
∼20 >500 15–30 ∼12 ∼160 ∼120 ∼80 ∼100 ∼35 ∼150 ∼30
This work [21] [24] [22] [25] [26] [27] [5] [28] [29] [29]
Fig. 2. (a) Push-in curves of Cf/Mu composites; (b) Morphologies of the pushed fiber after the push-in tests.
bon fiber. In addition, the diffraction pattern of the mullite matrix evidences a typical orthogonal structure.
displacement rate was 30 nm/s. The applied load and fiber displacement were continuously monitored during the tests. Standard shear-lag model was employed to derive the interface shear stress from the obtained load-displacement curves [20]. SENB test was performed following the ASTME 399-74 standard, to quantify the fracture toughness of the composites. The sample dimensions were ∼40 mm in length, ∼3 mm in width and ∼6 mm in thickness. The notch, with ∼3 mm in length and ∼0.3 mm in width, was introduced by metal-bonded diamond cutting in the thickness direction. The load was also applied in the same direction, with a span of 30 mm, to induce Mode I cracking of the notch. The crack propagation paths of the composites during SENB tests were observed by the local surface strain profile monitored by DIC technique. The strain evolution of the monitored region was obtained with a sampling rate of one image per second. The data was postprocessed by commercial ARAMIS software to analyze the strain field evolution during testing process. TPB tests were performed at 25, 600 and 1200 ◦ C in vacuum to study the temperature dependent fracture strength of the composite. The crosshead speed was 0.5 mm/min and the test span was 40 mm. The tested samples were machined to a thickness of ∼3 mm by surface grinder, a width of ∼4 mm and a length of ∼50 mm using a metal-bonded diamond cutting blade. The warp direction was parallel to the sample length direction.
3.2. Interfacial shear property
3. Results and discussion 3.1. Interface microstructure The as-received T300 carbon fibers were grooved in surface, with average diameters of ∼7 m (Fig. 1a). They were distributed randomly in the dense mullite matrix after the sol-gel process (Fig. 1b). The density of the as-fabricated Cf/Mu composite was 2.20 g cm−3 , and the fiber volume fraction was ∼43%. The fiber/matrix interfacial microstructure was characterized in detail by a combination of bright-field TEM (BF-TEM), high-resolution TEM (HR-TEM) and selected area diffraction (SAD) patterns, as is shown in Fig. 1(c)–(f). In BF-TEM mode, the fiber, matrix and interface can be distinctly distinguished. HR-TEM observations displayed an incoherent fiber/matrix interface microstructure, in which the (120) crystallized planes of the mullite matrix were parallel to the interface. No evidence of the interfacial reaction was observed, which is expected because the reported carbo-thermal reduction reaction between carbon and mullite is occurred at a temperature above 1800 ◦ C [13]. The SAD pattern on carbon fibers showed apparent (0001), etc. diffraction rings. The inhomogeneous contrast of the rings indicates a preferential texture of the car-
The interfacial shear properties of Cf/Mu composites were studied by fiber push-in method, employing a 3 m flat punch to generate pure shear stress along the interface, without cracking the carbon fiber. Note this method is different from what is known nanoindentation in which a typical sharp nanoindenter is employed to measure the Young’s modulus, hardness of the indented material. Fig. 2(a) shows the typical push-in force-displacement curves of the composites. The curves were reproducible, with initial linear response (stiffness labeled S0 ) until a transition to nonlinear at a critical force, Pc , which represents the debonding of the fibers from the matrix. As is evidenced in Fig. 2(b), after the push-in tests, the carbon fiber was completely debonded from the mullite matrix, without cracking the fiber or matrix. Based on the analytic shear lag model [20], the interfacial shear strength, , can be derived by: =
S0 Pc 22 r 3 Ef
(1)
where r is the average radius of the carbon fiber, and is ∼3.5 m (Fig. 1a). Ef is the longitudinal elastic modulus of the T300 carbon fiber, with a value of ∼210 GPa. Based on Eq. (1), the interfacial shear strength of the Cf/Mu composite was finally given 25 ± 5 MPa. Table 1 reviews the interfacial shear strength of some typical CMCs in literatures. Generally, the interfacial shear strength of most CMCs, e.g. SiCf/SiC, Cf/SiC, Al2 O3f /Al2 O3 etc., without the introduction of compliant interphase are in the orders of several hundred MPa (SiCf/SiC without interphase was about 700–800 MPa [21]), strong enough to promote crack penetration and lead to catastrophic brittle failures. However, this limitation can be overcome by the introduction of compliant interphase. E.g., the introduction of PyC interphase can lead to an interfacial shear strength of SiCf/SiC composites dropped from >500 MPa to ∼12 MPa [22]. Nevertheless, the interfacial shear strength of the interphase-free Cf/Mu composite is almost in the same order of that of the CMCs with compliant interphase. The formation mechanisms of the weak interface could be attributed to: firstly, no chemical reactions of the composite components were occurred (Fig. 1c), so the fiber/matrix interface was physically or mechanically controlled; secondly, due to the relatively large thermal expansion coefficient (␣) mismatch of the carbon fiber and mullite matrix (␣ for carbon fiber is ∼−0.4*10−6 /◦ C [20], and that of mullet ∼5*10−6 /◦ C for mullite [23]), large residual shear stress could be generated, which could further weaken the interface [14].
Please cite this article in press as: L.W. Yang, et al., Weak interface dominated high temperature fracture strength of carbon fiber reinforced mullite matrix composites, J Eur Ceram Soc (2017), http://dx.doi.org/10.1016/j.jeurceramsoc.2017.03.033
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3.3. Composite fracture toughness
3.4. High temperature properties
The weak fiber/matrix interface is expected to trigger the toughening mechanisms like crack deflection, crack arrest, etc., thus leading to enhancement of the fracture toughness of the composite, which was measured by the SENB test. Fig. 3(a) shows the SENB force-displacement curve of the Cf/Mu composite, showing a ductile response. Based on the standard ASTME 399-74, the √ composite fracture toughness was given 12.3 ± 1.2 MPa m. Compared to those of the as-received carbon fiber and mullite ceramic √ (typically ∼1 and ∼2 MPa m, respectively [30,31]), the dramatic enhancement of the fracture toughness in Cf/Mu composite should be mainly a consequence of the weak fiber/matrix interfacial interaction that triggers fully the toughening mechanisms, thus leading to large fiber pull-out length of the composite after the SENB test, shown in Fig. 3(b). The accumulated strain contours during composite fracturing (insets in Fig. 3a) further evidence a macro-scaled crack deflection during the fracture process, mainly due to the contribution of weak fiber/matrix interface.
Weak interface dominates the toughening behaviors of the Cf/Mu composites, and this effect was extended to high temperatures by performing high temperature TPB tests in a wide temperature range from 25 to 1200 ◦ C, and in vacuum to eliminate other environmental factors (composite oxidation, etc.). Fig. 4(a) shows the representative stress-displacement curves of TPB tests at 25, 600 and 1200 ◦ C. All curves exhibited “metal-like” responses, with elastic deformation up to a critical “yield” stress, followed by the energy dissipations including strain hardening and softening processes. The bending modulus of the composite (derived from the linear part of the curve) at 25 ◦ C, 600 ◦ C and 1200 ◦ C were 41.4(±3.8) GPa, 39.9(±2.5) GPa and 44.3(±3.4) GPa, respectively, which were independent of the testing temperature. This indicates a stable microstructure of the Cf/Mu composite in the temperature range considered. The fracture strength (determined as the maximum stress in the TPB curve) of the Cf/Mu composites can be an index to quantify the contribution of the weak interface to the toughening properties of the composites. At 25 ◦ C, the fracture stress of the composite was as high as ∼200 MPa, which is beyond that of most mullite matrix composites due to the weak interfacial interactions [32,33]. When increasing testing temperature up
Fig. 3. (a) SENB force-displacement curve for the Cf/Mu composite; (b) Fracture surface morphology after the SENB test. Inserts in (a) is the strain countours of the composite at different displacements.
Fig. 4. (a) TPB stress versus displacement of Cf/Mu composite at 25 ◦ C, 600 ◦ C and 1200 ◦ C; (b) The fracture surfaces of the composite after TPB test at 25 ◦ C and 1200 ◦ C.
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to 1200 ◦ C, an interesting upward tendency of the fracture stress versus temperature was observed, and at 1200 ◦ C, the composite’s fracture stress increased up to ∼280 MPa. The fracture surfaces of the composites after TPB tests at different temperatures all displayed long fiber pull-out features, as representatively shown in Fig. 4(b), which evidence that the toughening/failure mechanism of the composite is mainly triggered by the debonding of the interface, which is independent of the testing temperature. Therefore, a weak interface dominated high temperature fracture behavior can be concluded in the typical Cf/Mu composite. As for CMCs, the macro fracture strength is a function of localized properties of matrix, fiber and interface in a specific composite structure. Since mullite ceramics and carbon fibers could be stabilized in mechanical properties until extremely high temperatures (generally beyond 1200 ◦ C) [34,35], the evolution of the fracture strength of the Cf/Mu composite with testing temperature should be mainly a consequence of the temperature dependent interfacial properties. As discussed above, the weak interface in the as-fabricated Cf/Mu composite is mainly attributed to the chemical sharp interfacial microstructure (Fig. 2) and the residual shear thermal stress at the interface. While the interface microstructure is expected to be stabilized until 1800 ◦ C [13], the residual thermal stress at the interface could be released at elevated temperature. Accordingly, the coaction of the two factors yields an enhancement in the interfacial shear properties of the Cf/Mu composites at elevated temperatures. Therefore, the stronger behavior of the Cf/Mu composite at elevated temperatures is mainly a consequence of the enhanced interfacial property that favors better load transfer between the carbon fiber and mullite matrix. Note the strengthening of the fiber-matrix interface could decrease the toughness of the composite to some extent, due that the fracture energy is more difficult to be dissipated at the stronger interface. Nevertheless, the weak fiber/matrix interface not only favors the crack arrest/deflections to toughen the Cf/Mu composite, but also leads to an enhanced high temperature fracture strength of the composite. This specialty could facilitate Cf/Mu composites in high temperature structural applications.
4. Conclusions A novel carbon fiber reinforced mullite matrix composite was fabricated in this work by a conventional sol-gel process. The fiber/matrix interface was chemically free, and the interfacial shear strength, measured by fiber push-in tests, was ∼25 MPa, which is comparable to that of the typical ceramic matrix composites with traditional interphase. The weak interfacial interactions can fully trigger the toughening mechanisms, e.g. crack deflection, crack arrest, etc., of the composites, thus leading to a significant enhancement of the macro fracture toughness, which was ∼12.5 √ MPa m, measured by single edged notch beam method. Due to the weak interface interaction and large thermal expansion mismatch between the carbon fiber and mullite matrix, the high temperature fracture strength in the temperature range from 25 ◦ C and 1200 ◦ C was enhanced, and this is attributed to the enhancement of the interfacial property at elevated temperatures that favors better load transfers between composite constituents.
Acknowledgement Lingwei Yang greatly appreciates the financial support from the China Scholarship Council (grant number: 201306110007).
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Short communication
Weak interface dominated high temperature fracture strength of carbon fiber reinforced mullite matrix composites L.W. Yang b,∗ , H.T. Liu a,∗ , R. Jiang a , X. Sun a , W.G. Mao c , H.F. Cheng a , J.M. Molina-Aldareguía b a
Science and Technology on Advanced Ceramic Fibers and Composites Laboratory, National University of Defense Technology, Changsha 410073, China IMDEA Materials Institute, c/Eric Kandel 2, Getafe 28906, Madrid, Spain c School of Materials Science and Engineering, Xiangtan University, Xiangtan 411105, China b
a r t i c l e
i n f o
Article history: Received 23 October 2016 Received in revised form 14 March 2017 Accepted 17 March 2017 Available online xxx Keywords: A. Ceramic matrix composites B. Interfaces C. Fracture mechanics/toughness D. Strength
a b s t r a c t Weak fiber/matrix interface dominates the toughening properties of ceramic matrix composites. This paper reports a novel sol-gel fabricated carbon fiber reinforced mullite matrix composite, in which the fiber/matrix interface was inherently weak in shear properties (∼25 MPa), measured in-situ by fiber push-in tests. The interface microstructure was chemically sharp, characterized by transmission electron microscopy. The outcome of the weak interface was the full trigger of the toughening mechanisms like crack deflection, etc., leading to significant enhancement of the fracture toughness of the composite √ (∼12 MPa m), measured by single edged notch beam method. Finally, due to the weak fiber/matrix interface and large thermal expansion mismatch of the fiber and matrix, the high temperature fracture strength was enhanced in the temperature range from 25 to 1200 ◦ C, which is attributed to the enhancement of the interfacial property at elevated temperatures that favors better load transfers between composite constituents. © 2017 Published by Elsevier Ltd.
1. Introduction Ceramic matrix composites reinforced by continuous high strength fibers are promising materials with excellent thermal mechanical properties, e.g. high strength, improved toughness, low density, etc. [1,2]. The toughness of CMCs is still under study due to the concerns of the catastrophic brittle failures of CMCs in response to external stresses. Various methods have been developed to enhance the fracture resistance of the CMCs by tailoring the fiber/matrix interface. Among them, one common method is to introduce a compliant (ex-situ) interphase between the fiber and matrix by chemical vapor deposition, liquid phase processes etc., which are time-consuming and expensive [3,4]. Recently, some insitu methods have been proposed in carbon fiber reinforced SiC matrix (Cf/SiC) composites to fabricate the interphase directly by fiber heat treatment or by matrix fabrication [5,6]. These in-situ formed interphases have shown many advantages over the typical ex-situ interphases. Firstly, no precursors would be needed, avoiding the mechanical degradation of the carbon fibers; secondly, the
∗ Corresponding authors. E-mail addresses: [email protected] (L.W. Yang), [email protected] (H.T. Liu).
fabrication processes would be much easier and cheaper than most ex-situ processes; and finally, the in-situ interphase could be grown homogeneously, especially in the interior of three-dimensional fiber preforms, that are difficult to access by ex-situ coating methods. Despite that, an even more efficient, but challenging way to enhance the fracture resistance of CMCs is to introduce weak interface instead of compliant interphase directly during the composite fabrication process. However, the potential interfacial reactions between the fibers and the matrix during high temperature fabrications process would bring chemically controlled interfaces in many typical CMCs, thus leading to pretty strong interfacial interactions [7–10]. For example, in those SiC matrix composites that are fabricated by precursor infiltration and pyrolysis (PIP) of polycarbosilane (PCS), the Si H groups in PCS can be easily embedded by the O groups on fiber surfaces at high temperatures, forming O Si O like interphase that can bond the fiber/matrix interface tightly [9,11,12]. Therefore, proper choices of matrix and fiber materials to prevent potential interface reactions can promisingly enable the formation of weak interfaces in CMCs, without the introduction of ex-situ or in-situ interphases. As is known that mullite is a typical ceramic that can be hardly reacted with carbon materials due to the high carbo-thermal reduction reaction temperature (1800 ◦ C) [13], which could bring
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Fig. 1. (a) Surface morphology of the as-received T300 carbon fiber; (b) Polished cross-section of Cf/Mu; (c) BF-TEM of a TEM lamellae at the interface zone in Cf/Mu; (d) HR-TEM of the interface zone; (e) SAD pattern of the carbon fiber; (f) SAD pattern of the mullite matrix.
an inherently chemical sharp mullite/carbon interface. To take full advantage of this peculiarity, carbon fiber reinforced mullite matrix (Cf/Mu) composites have been developed decades ago [14]. Despite this, the interfacial properties in the Cf/Mu composites and their dependence on the interfacial microstructure and the macro fracture resistance of the Cf/Mu composites, especially at high temperatures, are still not fully explored. This information is critical not only from the fundamental viewpoint, but from the engineering perspectives since most CMCs are served in harsh environments. For the sake of filling the research gap, a novel Cf/Mu composite was fabricated in this work by sol-gel process. The fiber-matrix interface was characterized by transmission electron microscopy (TEM). The shear strength of the fiber-matrix interface was measured insitu by fiber push-in method. The mechanical fracture toughness of the composites was measured by single edged notch beam (SENB) test, and a digital image correction (DIC) method was employed to track the crack propagations to study the toughening mechanisms. Finally, the effect of the resulting interfacial property on the high temperature fracture strength of the Cf/Mu composite was studied by high temperature three-point bend (TPB) techniques in a wide temperature range from 25 to 1200 ◦ C in vaccum. 2. Experimental procedures PAN based 3 dimensional (3D) 3K-T300 carbon fiber fabrics were employed as the reinforcement. The fabric braid angle was 22◦
and the fiber volume fraction was ∼43%. The Diphasic mullite sols (AS12, from Suzhou Nanodispersions Ltd, China) prepared by comixing of silica and alumina sols were used as precursor of the mullite matrix. 3D Cf/Mu composites were fabricated by sol-gel process [15]. This method is advantageous over those traditional densification methods (slurry infiltration etc.) to achieve a denser mullite matrix at a comparably lower temperature, which is important to minimize the thermal-erosions and corrosions of the carbon fibers during matrix fabrication process [16–19]. The details of the sol-gel process are as follows: firstly, the mullite sols were infiltrated into carbon fiber preforms in a vacuum apparatus at room temperature for 4 h. After that, the infiltrated preform was sintered in Ar at 1200 ◦ C for 1 h after dried at 80 ◦ C for 12 h to gelate the mullite sols. Finally, the process was repeated for at least 14 times to obtain dense Cf/Mu composites with the relative composite weight increase less than 1% during subsequent sol-gel cycles. The morphologies of carbon fibers and Cf/Mu composites were observed by scanning electron microscope (SEM, Hitachi FEG S4800). The microstructure of the fiber/matrix interface zone was characterized by TEM (JEOL JEM 2100) with 200 kV accelerating voltage. The electron transparent TEM lamella with <100 nm in thickness was milled in a Helios 600i focused ion beam (FIB) system. The composite interface shear strength was evaluated by fiber push-in tests within a TI 950 TriboindenterTM nanoindentation system using a flat punch of 3 m in diameter. More than 8 push-in tests were carried out on selected fibers until fiber debonding. The
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Table 1 Review of the interfacial shear strength in the typical CMCs. (CVI: chemical vapor infiltration; PIP: precursor infiltration pyrolysis). Composites
Processing
Method
(MPa)
Ref
Cf/Mu SiCf/SiC SiCf/BN/SiC SiC/PyC/SiC SiCf/Mu Muf/ZrO2 /Al2 O3 Al2 O3 /LaPO4 /Al2 O3 Cf/SiC Cf/C/SiC Muf/SiC Muf/C/SiC
Sol-gel CVI CVI CVI Hot Press Cold Press + Sintering Sintering PIP CVI PIP PIP
Push in Push in Push out Push out Pull out Push in Push out Push in Push out Push in Push in
∼20 >500 15–30 ∼12 ∼160 ∼120 ∼80 ∼100 ∼35 ∼150 ∼30
This work [21] [24] [22] [25] [26] [27] [5] [28] [29] [29]
Fig. 2. (a) Push-in curves of Cf/Mu composites; (b) Morphologies of the pushed fiber after the push-in tests.
bon fiber. In addition, the diffraction pattern of the mullite matrix evidences a typical orthogonal structure.
displacement rate was 30 nm/s. The applied load and fiber displacement were continuously monitored during the tests. Standard shear-lag model was employed to derive the interface shear stress from the obtained load-displacement curves [20]. SENB test was performed following the ASTME 399-74 standard, to quantify the fracture toughness of the composites. The sample dimensions were ∼40 mm in length, ∼3 mm in width and ∼6 mm in thickness. The notch, with ∼3 mm in length and ∼0.3 mm in width, was introduced by metal-bonded diamond cutting in the thickness direction. The load was also applied in the same direction, with a span of 30 mm, to induce Mode I cracking of the notch. The crack propagation paths of the composites during SENB tests were observed by the local surface strain profile monitored by DIC technique. The strain evolution of the monitored region was obtained with a sampling rate of one image per second. The data was postprocessed by commercial ARAMIS software to analyze the strain field evolution during testing process. TPB tests were performed at 25, 600 and 1200 ◦ C in vacuum to study the temperature dependent fracture strength of the composite. The crosshead speed was 0.5 mm/min and the test span was 40 mm. The tested samples were machined to a thickness of ∼3 mm by surface grinder, a width of ∼4 mm and a length of ∼50 mm using a metal-bonded diamond cutting blade. The warp direction was parallel to the sample length direction.
3.2. Interfacial shear property
3. Results and discussion 3.1. Interface microstructure The as-received T300 carbon fibers were grooved in surface, with average diameters of ∼7 m (Fig. 1a). They were distributed randomly in the dense mullite matrix after the sol-gel process (Fig. 1b). The density of the as-fabricated Cf/Mu composite was 2.20 g cm−3 , and the fiber volume fraction was ∼43%. The fiber/matrix interfacial microstructure was characterized in detail by a combination of bright-field TEM (BF-TEM), high-resolution TEM (HR-TEM) and selected area diffraction (SAD) patterns, as is shown in Fig. 1(c)–(f). In BF-TEM mode, the fiber, matrix and interface can be distinctly distinguished. HR-TEM observations displayed an incoherent fiber/matrix interface microstructure, in which the (120) crystallized planes of the mullite matrix were parallel to the interface. No evidence of the interfacial reaction was observed, which is expected because the reported carbo-thermal reduction reaction between carbon and mullite is occurred at a temperature above 1800 ◦ C [13]. The SAD pattern on carbon fibers showed apparent (0001), etc. diffraction rings. The inhomogeneous contrast of the rings indicates a preferential texture of the car-
The interfacial shear properties of Cf/Mu composites were studied by fiber push-in method, employing a 3 m flat punch to generate pure shear stress along the interface, without cracking the carbon fiber. Note this method is different from what is known nanoindentation in which a typical sharp nanoindenter is employed to measure the Young’s modulus, hardness of the indented material. Fig. 2(a) shows the typical push-in force-displacement curves of the composites. The curves were reproducible, with initial linear response (stiffness labeled S0 ) until a transition to nonlinear at a critical force, Pc , which represents the debonding of the fibers from the matrix. As is evidenced in Fig. 2(b), after the push-in tests, the carbon fiber was completely debonded from the mullite matrix, without cracking the fiber or matrix. Based on the analytic shear lag model [20], the interfacial shear strength, , can be derived by: =
S0 Pc 22 r 3 Ef
(1)
where r is the average radius of the carbon fiber, and is ∼3.5 m (Fig. 1a). Ef is the longitudinal elastic modulus of the T300 carbon fiber, with a value of ∼210 GPa. Based on Eq. (1), the interfacial shear strength of the Cf/Mu composite was finally given 25 ± 5 MPa. Table 1 reviews the interfacial shear strength of some typical CMCs in literatures. Generally, the interfacial shear strength of most CMCs, e.g. SiCf/SiC, Cf/SiC, Al2 O3f /Al2 O3 etc., without the introduction of compliant interphase are in the orders of several hundred MPa (SiCf/SiC without interphase was about 700–800 MPa [21]), strong enough to promote crack penetration and lead to catastrophic brittle failures. However, this limitation can be overcome by the introduction of compliant interphase. E.g., the introduction of PyC interphase can lead to an interfacial shear strength of SiCf/SiC composites dropped from >500 MPa to ∼12 MPa [22]. Nevertheless, the interfacial shear strength of the interphase-free Cf/Mu composite is almost in the same order of that of the CMCs with compliant interphase. The formation mechanisms of the weak interface could be attributed to: firstly, no chemical reactions of the composite components were occurred (Fig. 1c), so the fiber/matrix interface was physically or mechanically controlled; secondly, due to the relatively large thermal expansion coefficient (␣) mismatch of the carbon fiber and mullite matrix (␣ for carbon fiber is ∼−0.4*10−6 /◦ C [20], and that of mullet ∼5*10−6 /◦ C for mullite [23]), large residual shear stress could be generated, which could further weaken the interface [14].
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3.3. Composite fracture toughness
3.4. High temperature properties
The weak fiber/matrix interface is expected to trigger the toughening mechanisms like crack deflection, crack arrest, etc., thus leading to enhancement of the fracture toughness of the composite, which was measured by the SENB test. Fig. 3(a) shows the SENB force-displacement curve of the Cf/Mu composite, showing a ductile response. Based on the standard ASTME 399-74, the √ composite fracture toughness was given 12.3 ± 1.2 MPa m. Compared to those of the as-received carbon fiber and mullite ceramic √ (typically ∼1 and ∼2 MPa m, respectively [30,31]), the dramatic enhancement of the fracture toughness in Cf/Mu composite should be mainly a consequence of the weak fiber/matrix interfacial interaction that triggers fully the toughening mechanisms, thus leading to large fiber pull-out length of the composite after the SENB test, shown in Fig. 3(b). The accumulated strain contours during composite fracturing (insets in Fig. 3a) further evidence a macro-scaled crack deflection during the fracture process, mainly due to the contribution of weak fiber/matrix interface.
Weak interface dominates the toughening behaviors of the Cf/Mu composites, and this effect was extended to high temperatures by performing high temperature TPB tests in a wide temperature range from 25 to 1200 ◦ C, and in vacuum to eliminate other environmental factors (composite oxidation, etc.). Fig. 4(a) shows the representative stress-displacement curves of TPB tests at 25, 600 and 1200 ◦ C. All curves exhibited “metal-like” responses, with elastic deformation up to a critical “yield” stress, followed by the energy dissipations including strain hardening and softening processes. The bending modulus of the composite (derived from the linear part of the curve) at 25 ◦ C, 600 ◦ C and 1200 ◦ C were 41.4(±3.8) GPa, 39.9(±2.5) GPa and 44.3(±3.4) GPa, respectively, which were independent of the testing temperature. This indicates a stable microstructure of the Cf/Mu composite in the temperature range considered. The fracture strength (determined as the maximum stress in the TPB curve) of the Cf/Mu composites can be an index to quantify the contribution of the weak interface to the toughening properties of the composites. At 25 ◦ C, the fracture stress of the composite was as high as ∼200 MPa, which is beyond that of most mullite matrix composites due to the weak interfacial interactions [32,33]. When increasing testing temperature up
Fig. 3. (a) SENB force-displacement curve for the Cf/Mu composite; (b) Fracture surface morphology after the SENB test. Inserts in (a) is the strain countours of the composite at different displacements.
Fig. 4. (a) TPB stress versus displacement of Cf/Mu composite at 25 ◦ C, 600 ◦ C and 1200 ◦ C; (b) The fracture surfaces of the composite after TPB test at 25 ◦ C and 1200 ◦ C.
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to 1200 ◦ C, an interesting upward tendency of the fracture stress versus temperature was observed, and at 1200 ◦ C, the composite’s fracture stress increased up to ∼280 MPa. The fracture surfaces of the composites after TPB tests at different temperatures all displayed long fiber pull-out features, as representatively shown in Fig. 4(b), which evidence that the toughening/failure mechanism of the composite is mainly triggered by the debonding of the interface, which is independent of the testing temperature. Therefore, a weak interface dominated high temperature fracture behavior can be concluded in the typical Cf/Mu composite. As for CMCs, the macro fracture strength is a function of localized properties of matrix, fiber and interface in a specific composite structure. Since mullite ceramics and carbon fibers could be stabilized in mechanical properties until extremely high temperatures (generally beyond 1200 ◦ C) [34,35], the evolution of the fracture strength of the Cf/Mu composite with testing temperature should be mainly a consequence of the temperature dependent interfacial properties. As discussed above, the weak interface in the as-fabricated Cf/Mu composite is mainly attributed to the chemical sharp interfacial microstructure (Fig. 2) and the residual shear thermal stress at the interface. While the interface microstructure is expected to be stabilized until 1800 ◦ C [13], the residual thermal stress at the interface could be released at elevated temperature. Accordingly, the coaction of the two factors yields an enhancement in the interfacial shear properties of the Cf/Mu composites at elevated temperatures. Therefore, the stronger behavior of the Cf/Mu composite at elevated temperatures is mainly a consequence of the enhanced interfacial property that favors better load transfer between the carbon fiber and mullite matrix. Note the strengthening of the fiber-matrix interface could decrease the toughness of the composite to some extent, due that the fracture energy is more difficult to be dissipated at the stronger interface. Nevertheless, the weak fiber/matrix interface not only favors the crack arrest/deflections to toughen the Cf/Mu composite, but also leads to an enhanced high temperature fracture strength of the composite. This specialty could facilitate Cf/Mu composites in high temperature structural applications.
4. Conclusions A novel carbon fiber reinforced mullite matrix composite was fabricated in this work by a conventional sol-gel process. The fiber/matrix interface was chemically free, and the interfacial shear strength, measured by fiber push-in tests, was ∼25 MPa, which is comparable to that of the typical ceramic matrix composites with traditional interphase. The weak interfacial interactions can fully trigger the toughening mechanisms, e.g. crack deflection, crack arrest, etc., of the composites, thus leading to a significant enhancement of the macro fracture toughness, which was ∼12.5 √ MPa m, measured by single edged notch beam method. Due to the weak interface interaction and large thermal expansion mismatch between the carbon fiber and mullite matrix, the high temperature fracture strength in the temperature range from 25 ◦ C and 1200 ◦ C was enhanced, and this is attributed to the enhancement of the interfacial property at elevated temperatures that favors better load transfers between composite constituents.
Acknowledgement Lingwei Yang greatly appreciates the financial support from the China Scholarship Council (grant number: 201306110007).
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