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carbon composites by rod-like SiOC ceramic
Diamond & Related Materials 102 (2020) 107673
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Optimizing PyC matrix interface to improve mechanical properties of carbon/carbon composites by rod-like SiOC ceramic
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T. Fenga,b, , H.J. Lina,b, H.J. Lia, , S.F. Wenb, M.D. Tonga ⁎
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C/C Composites Technology Research Center, Northwestern Polytechnical University, Xi'an 710072, China School of Mechanics, Civil Engineering & Architecture, Northwestern Polytechnical University, Xi'an, Shanxi 710072, PR China
ARTICLE INFO
ABSTRACT
Keywords: SiOC ceramic C/C composites Interface Mechanical properties
Rod-like SiOC ceramic in-situ reinforced carbon/carbon (C/C) composites have been synthesized in this work. The rod-like SiOC ceramic growth mechanism was investigated. By optimizing the interfaces, simultaneously, mechanical strength of C/Cs could be improved. Firstly, in order to protect and strengthen carbon fibers (CFs), about 200 nm carbon layer was deposited on the surface of CFs. Then SiOC ceramic were in-situ grown in the C/ Cs preform. After the densification through TGCVI, the bending and ILSS test revealed that these optimal designs including the carbon buffer layer and SiOC reinforcements between the fiber/matrix (F/M) and matrix/matrix (M/M) interfaces endowed C/Cs with improved bending strength by 118.85% and ILSS by 92.2%, respectively. Raman results also demonstrated that the defects in C/C composites could be significantly reduced by introducing SiOC ceramic and optimizing carbon layer.
1. Introduction C/C composites with the fiber braided-fabric as the reinforcements and pyrolytic carbon (PyC) as the matrix possess tremendous potential applications in the aerospace, automobile and marine industries because of their extremely strength, modulus, toughness, thermal stability, thermal shock resistance, thermal expansion, mechanical stability etc. [1–4]. As is frequently reported, mechanical properties are the major indicators to evaluate the worth of structural materials. As for C/ C composites, pyrolytic carbon grows around carbon fibers (CFts) layer by layer, which can cause interfaces from F/M (carbon fiber and PyC matrix) and M/M (PyC matrix layer and layer). Due to the existence of these interfaces, there will be defects in C/C composites, and these defects afford priority to crack initiation and propagation under stress and eventually lead to material failure. So there are still two main challenges limit its application [5–7]: (1) the poor adhesion among the carbon matrix (2) the low interfacial bonding strength. In order to enhance the bonding force within the matrix, many
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studies have reported C/C with different fillers used as secondary reinforcement. Li et al. [8] obtained high-performance C/C composites by in-situ introducing carbon nanofibers into C/C composites through chemical vapor infiltration (CVI) process. Lu et al. [9] introduced Si3N4 nanowires in order to reinforce C/C composites, and the out-ofplane compressive and interlaminar shear strength have been improved by 66.7% and 58% respectively. F/M interfaces, which influenced the crack propagation during failure, can serve as bridges and transfer the stress between carbon fibers and carbon matrix [10–12]. These interfaces would lead to failure, and also transfer oxidation in C/C composites [13,14]. So enhancing these interfaces could improve both the mechanical properties and the oxidation resistance of C/C composites. Specifically, the microstate of F/M interface has intrinsic correlation with the properties of C/C composites. In recent years, numerous investigations have been carried out to modify the F/M interface in order to obtain high performance C/C composites. Wu et al. [15] found that few defects and microvoids treatment on the surface of carbon fibers would improve the interface properties of C/C
Correspondence to: T. Feng, School of Mechanics, Civil Engineering & Architecture, Northwestern Polytechnical University, Xi'an, Shanxi 710072, PR China. Corresponding author. E-mail addresses: [email protected] (T. Feng), [email protected] (H.J. Li).
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https://doi.org/10.1016/j.diamond.2019.107673 Received 30 August 2019; Received in revised form 7 December 2019; Accepted 25 December 2019 Available online 27 December 2019 0925-9635/ © 2019 Published by Elsevier B.V.
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Fig. 1. Schematic of depositing PyC interface layer on CFs and followed by growth of rod-like SiOC by CVD to optimize F/M interface and strengthen PyC matrix.
composites. Feng et al. [16] obtained 3D C/C composites with an aligned in-situ grown CNTs on the surface of carbon cloths, which manifested the most notably improvements in out-of-plane and inplane compressive strength of 63% and 275% respectively. Chen et al. [17] used silicon carbide nanofibers to modify unidirectional carbon preform and obtained SiCNF-C/C composites. Results showed that the thermal conductivity and mechanical properties of C/C composites were both enhanced. The amorphous porous SiOC ceramics with excellent mechanical strength and extraordinary resistance to oxidation and corrosion could be considered as good candidates for structural applications at elevated temperatures [18–21]. Its unique porous structure can be infiltrated by liquid or gaseous substances to produce denser products for special applications. Therefore, it has been widely used as template and foam materials [22]. The uniformly distributed SiOC ceramics within the matrix materials can reduce shrinkages of the matrix and further improve mechanical properties of the matrix [23]. However, it rarely has been studied and researched reinforcing carbon matrix as a reinforcing. In order to improve comprehensive mechanical performance of the final C/C composites, in-situ rod-like SiOC ceramic modified C/C composites were prepared by a three-step technique involving sol-gel impregnation, chemical vapor infiltration (CVI) reaction and thermal gradient chemical vapor immersion (TGCVI) as our present work [24]. However, how to moderate the F/M interfacial bonding strength and supply effective reinforcements to the carbon matrix without degrading the fiber strength remain challenges get to achieve. In this article, a thin pyrocarbon (PyC) interface layer with thickness about 200 nm was deposited on the surface of CFts by CVI technique to optimize the F/M interfacial bonding. The PyC interface layers, else, provided several other advantages (1) strengthening the fibers to ensure the skeleton
structure of SiOC ceramic growth and CVI reaction; (2) protecting the fiber from corrosion during the impregnation process; (3) increasing the contact area with silica gel which can promote the efficiency of reaction between PyC and silica gel. Afterwards, rod-like SiOC ceramic were insitu grown in the C/C preform. After densifying, the in-situ SiOC ceramic enhanced the comprehensive mechanical performance of the C/Cs obtained. Three-point bending and interlaminar shearing (ILSS) testings were applied to examine the effect of the optimal carbon layer on the mechanical properties of C/Cs. The reinforcement mechanism has also been analyzed. 2. Experimental procedures 2.1. Preform preparation 2D needled carbon fiber felts (120 mm ∗ 80 mm ∗ 70 mm, Yixing Tianniao Co. Ltd., China) with a bulk density of 0.38–0.4 g/cm3 and fiber diameters of 6–8 μm were used in this work. In order to optimize and modify the interfaces within the final C/C composites, fiber felts were firstly deposited with a PyC interface layer by isothermal CVI technique, which was carried out at 1050 °C using flowing mixture of CH4 (80 L/h) and N2 (160 L/h) under the ambient pressure. The deposition time was 2 h. Afterwards PyC-CFts were soaked in the silicaprecursory sol which was made by tetraethoxysilane, ethanol, deionized water and HCl in a mass ratio of 2:1:0.5:0.005. After 4 days of gelation, by drying in the oven at 100 °C for 12 h, PyC-CFt/Sigel was obtained. Finally, they were heat-treated in the isothermal CVI (ICVI) furnace again. After 2 h CVI reaction also under conditions of temperature at 1050 °C, flowing mixture CH4 of 80 L/h and N2 of 160 L/h, rod-like SiOC ceramic in-situ grew within the carbon felts. The
Table 1 The general description of the samples. Sample name
Density
Description
Carbon felt (CFt) Carbon felt with pyrolytic carbon (PyC-CFt) Carbon felts/silica xerogel (CFt/Sigel) Carbon felt with pyrolitic carbon/silica xerogel (PyC-CFt/Sigel) C/C preform SiOC-C/C preform1 SiOC-C/C preform2 C/C composites SiOC- C/C1 SiOC- C/C2
0.38–0.40 g/cm3 0.43–0.46 g/cm3 0.57–0.6 g/cm3 0.60–0.63 g/cm3 0.6–0.65 g/cm3 0.7–0.75 g/cm3 0.7–0.75 g/cm3 1.72 g/cm3 1.78 g/cm3 1.80 g/cm3
2D needled carbon fiber felts carbon felt with prepared pyrolytic carbon about 200 nm CFt dispersed with silica xerogel PyC-CFt dispersed with silica xerogel CFt through CVI reaction progress CFt/Sigel PyC-CFt through CVI reaction progress PyC-CFt through CVI reaction progress SiOC-C/C preform 1 with densification process by thermal gradient CVI SiOC-C/C preform 1 with densification process by thermal gradient CVI SiOC-C/C preform 2 with densification process by thermal gradient CVI
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Fig. 2. SEM images of the surfaces of carbon fibers, (a) CFt and (b) PyC-CFt of X-Y axis, (c) CFt and (d) PyC-CFt of Z axis. Insert images are the higher magnifications.
synthesis of rod-like SiOC ceramic in carbon felts without PyC interface layer was also performed by isothermal CVI technique under identical growth conditions. The sample names of SiOC in carbon felts with and without PyC interface layer were SiOC-C/C preform 1and SiOC-C/C preform 2 respectively. The schematic of this manufacturing process is depicted in Fig. 1.
microscopy (FE-SEM, ZEISS-SUPRA 55, Germany), equipped with energy dispersive X-ray spectroscopy (EDS) to analyze the chemical element distribution of the materials. The interface morphology of F/M and M/M was analyzed by transmission electron microscope (TEM, FEI Tecnai F30G2, USA). The as-prepared composites were cut into 55 mm × 10 mm × 4 mm rectangles for and then polished bulk density and three-point-bend
2.2. Preform densification The densification was carried out by thermal gradient CVI technique with duration length of 150 h with natural gas as carbon source, nitrogen as dilute and carrier gas. During the TGCVI process, the temperature of the crucible was carefully controlled within the range of 920–1050 °C, and the flow rate of CH4 and N2 was kept at 80 L/h and 160 L/h respectively. The C/Cs containing both the PyC interface layer and SiOC were denoted as SiOC-C/C2, while the C/Cs containing only SiOC were denoted as SiOC-C/C1. The details of the samples studied in this paper were shown in Table 1. 2.3. Physical and mechanical properties The phase identification of the synthesized C/C, SiOC-C/C1 and SiOC-C/C2 were carried out by X-ray diffraction (XRD, PANalytical X'Pert Pro, Netherlands), which used Cu Kα radiation (λ = 1.5418°A). The structural and defects of C/C, SiOC-C/C1 and SiOC-C/C2 composites were analysed by Raman scattering spectrum (Raman, RENISHA Invia, England). The intensity ratio of the D peak (about 1360 cm−1) and G peak (about 1580 cm−1) named ID/IG was used to evaluate defects of materials. The surface morphology of the fiber and fractured samples were observed using a field emission scanning electron
Fig. 3. X-ray diffraction patterns for composites (a) C/C (b) SiOC-C/C1 (c) SiOC-C/C2. The insert is the spectra of SiOC ceramic.
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Fig. 4. SEM images: inside of PyC-CFt/Sigel (a), PyC-CFt/Sigel after 1050 °C heat treating without PyC (b) and SiOC-C/C preform 2 (c); EDS and element content of spot1 (d), spot2 (e) and spot3 (f) corresponding (a), (b) and (c) respectively. The inserts are images at different magnifications.
testing. The dimensions of the samples were 10 mm × 10 mm × 10 mm for the ILSS test. The Archimedes method was applied to measure the density of the as-prepared samples. The flexural strength by three-pointbend testing was conducted by an electronic universal testing machine (CMT 5304, Suns Co. China), with the across-head speed of 0.5 mm/min and a span of 25 mm at room temperature. And all of the testing, at least five samples were tested.
3. Results and discussion 3.1. Morphology and structural SEM images of the surfaces and cross sections from CFt and PyC-CFt were shown in Fig. 2. By comparing the surfaces of CFt and PyC-CFt, XY axis of the PyC-CFt surface is rougher than CFt, coated with PyC with thickness about 200 nm from the Z axis.
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treating without PyC and SiOC-C/C preform 2 (Fig. 4(a–c)) were captured using SEM. In Fig. 4(a), Si gel grains of different sizes uniformly dispersed in the CFts and attached to the fibers as the result of the dehydration of Si sol. The corresponding EDS spectrum (Fig. 4(d)) showed that the three main ingredients were Si, O and C, whose ratio of atoms number was about 1:2:5. SiOC-C/C preform 2 was prepared by heat-treating PyC-CFt/Sigel at 1050 °C for 2 h, but CH4 was not introduced in this process. As Fig. 4(b) shown, irregular bulk silica materials also attached the fibers, with some irregular round pores on it. The magnified image (top right corner) indicated that the silica materials presented the varying aperture size porous appearance. This morphology was similar to the Si-O-C materials which prepared through a sol-gel method by Liu et al. [30]. From EDS analysis (Fig. 4(e)), these porous materials also contained O, Si and C elements. Compared with PyC-CFt/Sigel, carbon content was decreased. O and Si with an atom ratio of 2:1 were major components. The SiOC ceramic was shown in SEM image of Fig. 4(c), where SiOC phase and fiber phase mixed together and accompanied with many microcracks. From the enlarged graph in Fig. 4(f), the internal was full of rod-like structural, which were composed of elemental Si, O and C revealed by EDS analysis. Due to CVI process, carbon from PyC matrix and SiOC accounted for a major fraction of the materials (76.99% of the total, Atomic). According to the above images, a hypothesis of the possible formation mechanism of the rod-like SiOC ceramic was hereby drawn. As Fig. 5 shown, firstly, with the gradually increasing temperature, Si gel gradually formed a porous (SiO)x ceramic, at the same time, the free carbon from CH4 decomposition during CVI enclosed around the material infiltrated into the microspores. Then, the interaction among C and SiO groups occurred, and eventually formed the amorphousstructured SiOC material. A similar mechanism analysis had also been discussed by Li et al. [31]. Raman spectroscopy is used to evaluate the degree of orderliness or randomness within the carbon materials, because quantitative analysis can be achieved by the intensity ratio of the D peak and G peak in Raman spectra [32]. All the spectra taken from the three composites exhibited two prominent peaks, which were the disorder-induced D mode (~1360 cm−1) and the tangential G mode (~1580 cm−1), as shown in Fig. 6. From the spectra of C/C composites, we can see the maximal intensity ratio of the D peak and G peak, which illustrated that there were many defects in the C/C composites. These defects may be from the interfaces and densification process. Due to the typical vibration modes in carbon based materials [33,34], the Raman spectrum of the SiOC materials [35] also showed the presence of two high intensity bands at about 1360 and 1580 cm−1 which corresponded D peak and G peak respectively. Therefore there was no obvious peak location shift among SiOC-C/C1 (b), SiOC-C/C2 (c) and C/C (a). But the ratio of the intensities of the D and G bands (ID/IG) were determined to be 1.35, 1.17 and 1.09, respectively. Comparing these ratio with C/C, SiOC-C/C1 and SiOC-C/C2, the results suggest that there were a better bond and fewer defects among PyC matrix. As can be seen in curve c in Fig. 6, SiOC-C/C2 had the lowest value of ID/IG, which may be contributed by the PyC interface layer among fibers and carbon matrixes. With the bond density within the interlayer increased, the density increase and the defects decreased [36,37]. In a nutshell, high density and few defects could lead to a significant increase in mechanical strength of composites.
Fig. 5. Schematic structure of SiOC material.
Fig. 6. Raman spectra for composites (a) C/C (b) SiOC-C/C1 (c) SiOC-C/C2.
Fig. 3 showed the XRD patterns of C/C (a), SiOC-C/C1 (b) and SiOCC/C2 (c), and the insert was the spectra of SiOC ceramic. From the patterns, one obvious diffraction peak at about 26° was observed, which was assigned to (002) diffraction plane of carbon fiber [25,26]. The spectra of the SiOC ceramics with a broad peak between 20° and 30° indicated an entirely amorphous structure, which demonstrated the presence of non-crystalline silica [27–29]. The overlapping peak between C and SiOC lead to no obvious difference among the three samples. Although C and SiOC phases coexist in the sample, the structure did not changed muchly. To investigate the growth mechanism of rod-like SiOC ceramic, cross-section images of PyC-CFt/Sigel, PyC-CFt/Sigel after 1050 °C heat
3.2. Physical and mechanical properties Fig. 7 showed the three-point bending test results of C/C, SiOC-C/
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Fig. 7. Three-point bending stress-strain curves (a) average bending strengths (b) and bending maximum loading (c) of the samples.
C1 and SiOC-C/C2. The distinctive behaviors in these samples were presented by the curves with different slopes in Fig. 7(a). Compared with C/C, the bending strengths of SiOC-C/C1 was up slightly, which means that rod-like SiOC ceramic can improve the bending strengths. But without optimizing carbon layer interface, the mechanical properties could not be improved obviously. SiOC-C/C2 showed the highest slope than the rest samples. This observation suggests that both introducing rod-like SiOC ceramic and optimizing interface by PyC can enhance the bending performance. We hereby speculated that fractures could gradually occur during the loading process. Firstly, some fibers fractured and failed, which would led to fine incline of the bending performance. It was obvious that introducing SiOC can increase the brittleness of SiOC-C/C1 and SiOC-C/C2 materials, which then will break under higher loading. In addition, among fiber bundles, SiOC functions as a strengthening agent, which could then fix the fiber bundles and keep them from separating. The bending strengths of SiOCC/C1 and SiOC-C/C2 composites have improved by nearly 37.22% and 118.85%, compared with that of C/C composites, and the bending maximum loading had also increased by 34.60% and 121.42%, respectively.
Fig. 8(a) showed the ILSS stress-strain curves obtained corresponding to C/C, SiOC-C/C1 and SiOC-C/C2 specimens. The curves displayed a clear the specimens' response behaviors. As seen by the difference in slopes, response curves of SiOC-C/C2 specimens showed a higher slope than those of the C/C and SiOC-C/C1 specimens. This observation suggests that the former composites were stiffer than the latter ones. Fig. 8(b) and (c) showed the average shear strengths and shear maximum loading of the samples. Apparently the shear strength (Fig. 8(b)) and shear maximum loading (Fig. 8(c)) of SiOC-C/C1 and SiOC-C/C2 were significantly increased after the in-situ growth of SiOC ceramic in the C/C composites, especially for the composites optimized by the PyC interlayer. Compared with C/C composites, the results indicated that the shear strengths of SiOC-C/C1 and SiOC-C/C2 have improved by 45.7% and 92.2%, respectively. Shear maximum loading of SiOC-C/C1 and SiOC-C/C2 also have improved by 23.44% and 57.46%, respectively. 3.3. Fracture behavior In order to further analyze the enhancement mechanism of C/C
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Fig. 8. ILSS stress-strain curves (a) average shear strengths (b) and shear maximum loading (c) of the samples.
composites by in-situ grown rod-like SiOC ceramic, the fracture surface of the bend tested samples was investigated. The results were shown in Fig. 9, which displayed fracture surface of C/C, and different-magnification SEM images of SiOC-reinforced C/C composites. When the flexural stress is loaded on the composite samples, the strength of the composites is mainly depended on the strength of CFs. Based on the data of Fig. 7 and SEM Fig. 9(a), the fracture process in C/C during bending test can be inferred as follows: when the bending stress was loaded on the C/C samples, destructive cracks originated somewhere within the matrix at the most critical flaws [38]], then propagated along the poor F/M interface, which led to some fibers being pulled out from the PyC. These fibers first were destroyed in failure, which substantially weakened the overall performance of the materials. As the loading increased, the cracks from PyC matrix expanded, and failure rates of fibers continued until thorough destruction disruption [16]. As for SiOC-C/C2, the fracture surface was flat and with nearly absented of fiber pullout (Fig. 9(b)). This fracture surface could be attributed to the strongly-enhanced cohesion in matrix and also the powerful mechanical interlocking at F/M interface, which contributed to a straight and total break rather than step-by-step. Additionally, from the magnification SEM, SiOC ceramic rods pulling out from PyC matrix (Fig. 9(c)) and bridging between fiber and PyC (Fig. 9(d)) could also be found.
Bridging could effectively inhibit cracks to extend. Because inhibited crack expansion and reinforced phase pullout require a great amount of fractured energy consumption during the failure process [15,39], this mechanism in turn increased the flexural strength and drove a gradual release of the flexural stress. The fracture toughness of SiOC-C/C2 also had a greatly improved compared with C/C. The corresponding EDS spectrums (Fig. 9(e) and (f)) from spot 1 (Fig. 9(a)) and spot 2 (Fig. 9(b)) proved the existence of SiOC. When the interlaminar shear stress is loaded on the composite samples, the shear strength is mainly depended on both the matrix cohesion and F/M interfacial bonding strength [38]. Fig. 10 presented the shearing fracture surface of three composites. Firstly, Fig. 10(a–c) showed fracture surfaces paralleled to the fibers. As seen in Fig. 10a, the fractured surface of C/C showed flat and few PyC existed on the surface of fibers (inset), which suggested a poor bonding strength between PyC matrix and Fibers. As for the SiOC-C/C1 (Fig. 10(b)), matrix delaminating from fibers was inhibited and abundant damaged SiOC can be clearly observed in the shearing fracture surface, indicating that the interfacial strength between fibers and PyC was stronger than C/C. This phenomenon also prevented crack expansion along the interface between fiber and PyC matrix. Thus this mechanism led to cracks deflection and fractured energy consumption. Especially for SiOC-C/C2
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Fig. 9. SEM images of the bending failure specimens from C/C (a) and SiOC-C/C2 (different-magnification, b, c and d), EDS spectrums from spot 1 (e) and spot 2 (f).
(Fig. 10(c)), because of optimized the interlayer of F/M, PyC tightly attached on the surface of fibers. Therefore, C/Cs the matrix cohesion is lower than F/M interfacial bonding strength. This observation provided a direct evidence that the carbon interface layer can significantly improve interfacial bonding strength between fibers and PyC matrix. It also could be testified by the Fig. 10(d–f) which showed the fracture
surfaces perpendicular to the fibers of C/C, SiOC-C/C1 and SiOC-C/C2 respectively. Compared with C/C, the adhesiveness of SiOC-C/C1 and SiOC-C/C2 were both improved, and SiOC-C/C2 was the best among them. At the same time, the cracks in the matrix became less obvious, and the defects in this material could be reduced drastically, which also was proved by Raman result.
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Fig. 10. (a) SEM images of the shearing fracture surfaces of the three composites: (a–c) parallel to the fibers and (d–f) perpendicular to the fibers (insets are the magnified images) of the C/C, SiOC-C/C1 and SiOC-C/C2 respectively.
TEM images of the F/M and M/M interfaces were reported in Fig. 11. A layered structure was observed in the SiOC-C/C2 composites, and around the fiber, PyC phase and SiOC particles can also be found. Already at low magnification it could be seen that cracking was induced at the F/M interface, owing to the external force during the sample preparation. But the crack stopped propagating when it encountered the SiOC particle. So it can be inferred that the SiOC can prevent crack growth from the failure occurring along the interfaces under mechanical loading. SiOC can strengthen the interfaces of the SiOC-C/C2 composites, which made the material have stronger mechanical
properties and not easily to be destroyed. Even a lot of cracked SiOC ceramic rods can be observed in the matrix fracture of SiOC-C/C2 (Fig. 12). So introducing rod-like SiOC ceramic into C/C composites could greatly improve both the matrix cohesion and F/M interfacial bonding strength. For SiOC-C/C2, on one hand, an excellent interface bonding can effectively inhibit the initiation and propagation of cracks along the F/M interface. On the other hand, SiOC ceramic rods' own strength and high matrix cohesion all consume the stress. It also explained why the increment in ILSS of SiOCC/C2 (that is 92.2%) was higher than that of SiOC-C/C1 (that is 45.7%).
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Fig. 11. Typical TEM micrographs of the F/M and M/M interfaces domain in the SiOC-C/C2 composites: (a) low magnification and (b) high magnification with the EDS spectrum.
Fig. 12. SEM images cross-section of failure specimens of SiOC-C/C2.
4. Conclusions
mechanical strength to replace traditional C/C in industries.
Rod-like SiOC ceramics were in-situ grown in the C/C composites by CVI reaction method using the porous properties of (SiO)x ceramic. In order to optimize F/M interface, strength and protect fibers, about 200 nm carbon layer has been deposited on the CFs surface at the beginning. Three-point bending and ILSS testing results showed that the above-mentioned material named SiOC-C/C2 have nearly 118.85% and 92.2% increase improvements compare to C/C composites in bending and shear strengths respectively. SEM morphologies of fracture surfaces of failure composites revealed that the synergistic effects of grown rodlike SiOC ceramic and carbon layer interface along with optimized F/M interfaces not only efficiently improved the matrix cohesion, but also the F/M interfacial bonding strength, which significantly decrease cracks generated and inhibited the crack propagation. This work opens up a possibility to produce SiOC-reinforced C/Cs with excellent
Author statement I have revised the manuscript of “Optimizing PyC matrix interface to improve mechanical properties of carbon/carbon composites by rodlike SiOC ceramic”. I have approved the final version to be published. All authors agree to publish this article, and all the dates are original. I certify that I did not submit it to other journals. Declaration of competing interest We declare that we have no financial and personal relationships with other people or organizations that can inappropriately influence
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our work. There is no professional or other personal interest of any nature or kind in any product, service or company that could be construed as influencing the position presented in, or the review of, the manuscript entitled” Optimizing PyC matrix interface to improve mechanical properties of carbon/carbon composites by rod-like SiOC ceramic”.
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Acknowledgements This work has been supported by the National Natural Science Foundation of China under Grant Nos. 51872237 and 51521061, and Natural Science Basic Research Plan in Shaanxi Province of China No. 2017JM5098. References [1] G.B. Zheng, H. Sano, Y. Uchiyama, A carbon nanotube-enhanced SiC coating for the oxidation protection of C/C composite materials, Compos Part B-Eng 42 (2011) 2158–2162. [2] J. Cheng, H.J. Li, S.Y. Zhang, L.Z. Xue, W.F. Luo, W. Li, Failure behavior investigation of a unidirectional carbon-carbon composite, Mater Design 55 (2014) 846–850. [3] Z.H. Hou, W. Yang, J.S. Li, R.Y. Luo, H.Z. Xu, Densification kinetics and mechanical properties of carbon/carbon composites reinforced with carbon nanotubes produced in situ, Carbon 99 (2016) 533–540. [4] X.R. Ren, H.J. Li, Q.G. Fu, K.Z. Li, Oxidation protective TaB2-SiC gradient coating to protect SiC-Si coated carbon/carbon composites against oxidation, Compos Part BEng 66 (2014) 174–179. [5] L. Li, H.J. Li, X.M. Yin, Y.H. Chu, X. Chen, Q.G. Fu, Oxidation protection and behavior of in-situ zirconium diboride-silicon carbide coating for carbon/carbon composites, J Alloy Compd 645 (2015) 164–170. [6] Q.G. Fu, B.Y. Tan, L. Zhuang, J.Y. Jing, Significant improvement of mechanical properties of carbon/carbon composites by in situ growth of SiC nanowires, Mat Sci Eng a-Struct 672 (2016) 121–128. [7] Y.J. Kwon, Y. Kim, H. Jeon, S. Cho, W. Lee, J.U. Lee, Graphene/carbon nanotube hybrid as a multi-functional interfacial reinforcement for carbon fiber-reinforced composites, Compos Part B-Eng 122 (2017) 23–30. [8] H.L. Li, H.J. Li, J.H. Lu, K.Z. Li, C. Sun, D.S. Zhang, Mechanical properties enhancement of carbon/carbon composites by in situ grown carbon nanofibers, Mat Sci Eng a-Struct 547 (2012) 138–141. [9] J.H. Lu, K.B. Guo, Q. Song, Y.Y. Li, L.L. Zhang, H.J. Li, In-situ synthesis silicon nitride nanowires in carbon fiber felts and their effect on the mechanical properties of carbon/carbon composites, Mater Design 99 (2016) 389–395. [10] X.F. Lu, P. Xiao, Short time oxidation behavior and residual mechanical properties of C/C composites modified by in situ grown carbon nanofibers, Ceram. Int. 40 (2014) 10705–10709. [11] Y. Swolfs, J. Shi, Y. Meerten, P. Hine, I. Ward, I. Verpoest, L. Gorbatikh, The importance of bonding in intralayer carbon fibre/self-reinforced polypropylene hybrid composites, Compos Part a-Appl S 76 (2015) 299–308. [12] M. Sharma, S.L. Gao, E. Mader, H. Sharma, L.Y. Wei, J. Bijwe, Carbon fiber surfaces and composite interphases, Compos. Sci. Technol. 102 (2014) 35–50. [13] C. Unterweger, J. Duchoslav, D. Stifter, C. Furst, Characterization of carbon fiber surfaces and their impact on the mechanical properties of short carbon fiber reinforced polypropylene composites, Compos. Sci. Technol. 108 (2015) 41–47. [14] X.F. Lu, P. Xiao, J. Chen, Y. Long, Oxidation behavior of C/C composites with the fibre/matrix interface modified by carbon nanotubes grown in situ at low temperature, Corros. Sci. 55 (2012) 20–25. [15] S. Wu, Y.Q. Liu, Y.C. Ge, L.P. Ran, K. Peng, M.Z. Yi, Surface structures of PAN-based carbon fibers and their influences on the interface formation and mechanical properties of carbon-carbon composites, Compos Part a-Appl S 90 (2016) 480–488. [16] L. Feng, K.Z. Li, Z.G. Zhao, H.J. Li, L.L. Zhang, J.H. Lu, Q. Song, Three-dimensional carbon/carbon composites with vertically aligned carbon nanotubes: providing direct and indirect reinforcements to the pyrocarbon matrix, Mater Design 92
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Optimizing PyC matrix interface to improve mechanical properties of carbon/carbon composites by rod-like SiOC ceramic
T
T. Fenga,b, , H.J. Lina,b, H.J. Lia, , S.F. Wenb, M.D. Tonga ⁎
a b
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C/C Composites Technology Research Center, Northwestern Polytechnical University, Xi'an 710072, China School of Mechanics, Civil Engineering & Architecture, Northwestern Polytechnical University, Xi'an, Shanxi 710072, PR China
ARTICLE INFO
ABSTRACT
Keywords: SiOC ceramic C/C composites Interface Mechanical properties
Rod-like SiOC ceramic in-situ reinforced carbon/carbon (C/C) composites have been synthesized in this work. The rod-like SiOC ceramic growth mechanism was investigated. By optimizing the interfaces, simultaneously, mechanical strength of C/Cs could be improved. Firstly, in order to protect and strengthen carbon fibers (CFs), about 200 nm carbon layer was deposited on the surface of CFs. Then SiOC ceramic were in-situ grown in the C/ Cs preform. After the densification through TGCVI, the bending and ILSS test revealed that these optimal designs including the carbon buffer layer and SiOC reinforcements between the fiber/matrix (F/M) and matrix/matrix (M/M) interfaces endowed C/Cs with improved bending strength by 118.85% and ILSS by 92.2%, respectively. Raman results also demonstrated that the defects in C/C composites could be significantly reduced by introducing SiOC ceramic and optimizing carbon layer.
1. Introduction C/C composites with the fiber braided-fabric as the reinforcements and pyrolytic carbon (PyC) as the matrix possess tremendous potential applications in the aerospace, automobile and marine industries because of their extremely strength, modulus, toughness, thermal stability, thermal shock resistance, thermal expansion, mechanical stability etc. [1–4]. As is frequently reported, mechanical properties are the major indicators to evaluate the worth of structural materials. As for C/ C composites, pyrolytic carbon grows around carbon fibers (CFts) layer by layer, which can cause interfaces from F/M (carbon fiber and PyC matrix) and M/M (PyC matrix layer and layer). Due to the existence of these interfaces, there will be defects in C/C composites, and these defects afford priority to crack initiation and propagation under stress and eventually lead to material failure. So there are still two main challenges limit its application [5–7]: (1) the poor adhesion among the carbon matrix (2) the low interfacial bonding strength. In order to enhance the bonding force within the matrix, many
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studies have reported C/C with different fillers used as secondary reinforcement. Li et al. [8] obtained high-performance C/C composites by in-situ introducing carbon nanofibers into C/C composites through chemical vapor infiltration (CVI) process. Lu et al. [9] introduced Si3N4 nanowires in order to reinforce C/C composites, and the out-ofplane compressive and interlaminar shear strength have been improved by 66.7% and 58% respectively. F/M interfaces, which influenced the crack propagation during failure, can serve as bridges and transfer the stress between carbon fibers and carbon matrix [10–12]. These interfaces would lead to failure, and also transfer oxidation in C/C composites [13,14]. So enhancing these interfaces could improve both the mechanical properties and the oxidation resistance of C/C composites. Specifically, the microstate of F/M interface has intrinsic correlation with the properties of C/C composites. In recent years, numerous investigations have been carried out to modify the F/M interface in order to obtain high performance C/C composites. Wu et al. [15] found that few defects and microvoids treatment on the surface of carbon fibers would improve the interface properties of C/C
Correspondence to: T. Feng, School of Mechanics, Civil Engineering & Architecture, Northwestern Polytechnical University, Xi'an, Shanxi 710072, PR China. Corresponding author. E-mail addresses: [email protected] (T. Feng), [email protected] (H.J. Li).
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https://doi.org/10.1016/j.diamond.2019.107673 Received 30 August 2019; Received in revised form 7 December 2019; Accepted 25 December 2019 Available online 27 December 2019 0925-9635/ © 2019 Published by Elsevier B.V.
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Fig. 1. Schematic of depositing PyC interface layer on CFs and followed by growth of rod-like SiOC by CVD to optimize F/M interface and strengthen PyC matrix.
composites. Feng et al. [16] obtained 3D C/C composites with an aligned in-situ grown CNTs on the surface of carbon cloths, which manifested the most notably improvements in out-of-plane and inplane compressive strength of 63% and 275% respectively. Chen et al. [17] used silicon carbide nanofibers to modify unidirectional carbon preform and obtained SiCNF-C/C composites. Results showed that the thermal conductivity and mechanical properties of C/C composites were both enhanced. The amorphous porous SiOC ceramics with excellent mechanical strength and extraordinary resistance to oxidation and corrosion could be considered as good candidates for structural applications at elevated temperatures [18–21]. Its unique porous structure can be infiltrated by liquid or gaseous substances to produce denser products for special applications. Therefore, it has been widely used as template and foam materials [22]. The uniformly distributed SiOC ceramics within the matrix materials can reduce shrinkages of the matrix and further improve mechanical properties of the matrix [23]. However, it rarely has been studied and researched reinforcing carbon matrix as a reinforcing. In order to improve comprehensive mechanical performance of the final C/C composites, in-situ rod-like SiOC ceramic modified C/C composites were prepared by a three-step technique involving sol-gel impregnation, chemical vapor infiltration (CVI) reaction and thermal gradient chemical vapor immersion (TGCVI) as our present work [24]. However, how to moderate the F/M interfacial bonding strength and supply effective reinforcements to the carbon matrix without degrading the fiber strength remain challenges get to achieve. In this article, a thin pyrocarbon (PyC) interface layer with thickness about 200 nm was deposited on the surface of CFts by CVI technique to optimize the F/M interfacial bonding. The PyC interface layers, else, provided several other advantages (1) strengthening the fibers to ensure the skeleton
structure of SiOC ceramic growth and CVI reaction; (2) protecting the fiber from corrosion during the impregnation process; (3) increasing the contact area with silica gel which can promote the efficiency of reaction between PyC and silica gel. Afterwards, rod-like SiOC ceramic were insitu grown in the C/C preform. After densifying, the in-situ SiOC ceramic enhanced the comprehensive mechanical performance of the C/Cs obtained. Three-point bending and interlaminar shearing (ILSS) testings were applied to examine the effect of the optimal carbon layer on the mechanical properties of C/Cs. The reinforcement mechanism has also been analyzed. 2. Experimental procedures 2.1. Preform preparation 2D needled carbon fiber felts (120 mm ∗ 80 mm ∗ 70 mm, Yixing Tianniao Co. Ltd., China) with a bulk density of 0.38–0.4 g/cm3 and fiber diameters of 6–8 μm were used in this work. In order to optimize and modify the interfaces within the final C/C composites, fiber felts were firstly deposited with a PyC interface layer by isothermal CVI technique, which was carried out at 1050 °C using flowing mixture of CH4 (80 L/h) and N2 (160 L/h) under the ambient pressure. The deposition time was 2 h. Afterwards PyC-CFts were soaked in the silicaprecursory sol which was made by tetraethoxysilane, ethanol, deionized water and HCl in a mass ratio of 2:1:0.5:0.005. After 4 days of gelation, by drying in the oven at 100 °C for 12 h, PyC-CFt/Sigel was obtained. Finally, they were heat-treated in the isothermal CVI (ICVI) furnace again. After 2 h CVI reaction also under conditions of temperature at 1050 °C, flowing mixture CH4 of 80 L/h and N2 of 160 L/h, rod-like SiOC ceramic in-situ grew within the carbon felts. The
Table 1 The general description of the samples. Sample name
Density
Description
Carbon felt (CFt) Carbon felt with pyrolytic carbon (PyC-CFt) Carbon felts/silica xerogel (CFt/Sigel) Carbon felt with pyrolitic carbon/silica xerogel (PyC-CFt/Sigel) C/C preform SiOC-C/C preform1 SiOC-C/C preform2 C/C composites SiOC- C/C1 SiOC- C/C2
0.38–0.40 g/cm3 0.43–0.46 g/cm3 0.57–0.6 g/cm3 0.60–0.63 g/cm3 0.6–0.65 g/cm3 0.7–0.75 g/cm3 0.7–0.75 g/cm3 1.72 g/cm3 1.78 g/cm3 1.80 g/cm3
2D needled carbon fiber felts carbon felt with prepared pyrolytic carbon about 200 nm CFt dispersed with silica xerogel PyC-CFt dispersed with silica xerogel CFt through CVI reaction progress CFt/Sigel PyC-CFt through CVI reaction progress PyC-CFt through CVI reaction progress SiOC-C/C preform 1 with densification process by thermal gradient CVI SiOC-C/C preform 1 with densification process by thermal gradient CVI SiOC-C/C preform 2 with densification process by thermal gradient CVI
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Fig. 2. SEM images of the surfaces of carbon fibers, (a) CFt and (b) PyC-CFt of X-Y axis, (c) CFt and (d) PyC-CFt of Z axis. Insert images are the higher magnifications.
synthesis of rod-like SiOC ceramic in carbon felts without PyC interface layer was also performed by isothermal CVI technique under identical growth conditions. The sample names of SiOC in carbon felts with and without PyC interface layer were SiOC-C/C preform 1and SiOC-C/C preform 2 respectively. The schematic of this manufacturing process is depicted in Fig. 1.
microscopy (FE-SEM, ZEISS-SUPRA 55, Germany), equipped with energy dispersive X-ray spectroscopy (EDS) to analyze the chemical element distribution of the materials. The interface morphology of F/M and M/M was analyzed by transmission electron microscope (TEM, FEI Tecnai F30G2, USA). The as-prepared composites were cut into 55 mm × 10 mm × 4 mm rectangles for and then polished bulk density and three-point-bend
2.2. Preform densification The densification was carried out by thermal gradient CVI technique with duration length of 150 h with natural gas as carbon source, nitrogen as dilute and carrier gas. During the TGCVI process, the temperature of the crucible was carefully controlled within the range of 920–1050 °C, and the flow rate of CH4 and N2 was kept at 80 L/h and 160 L/h respectively. The C/Cs containing both the PyC interface layer and SiOC were denoted as SiOC-C/C2, while the C/Cs containing only SiOC were denoted as SiOC-C/C1. The details of the samples studied in this paper were shown in Table 1. 2.3. Physical and mechanical properties The phase identification of the synthesized C/C, SiOC-C/C1 and SiOC-C/C2 were carried out by X-ray diffraction (XRD, PANalytical X'Pert Pro, Netherlands), which used Cu Kα radiation (λ = 1.5418°A). The structural and defects of C/C, SiOC-C/C1 and SiOC-C/C2 composites were analysed by Raman scattering spectrum (Raman, RENISHA Invia, England). The intensity ratio of the D peak (about 1360 cm−1) and G peak (about 1580 cm−1) named ID/IG was used to evaluate defects of materials. The surface morphology of the fiber and fractured samples were observed using a field emission scanning electron
Fig. 3. X-ray diffraction patterns for composites (a) C/C (b) SiOC-C/C1 (c) SiOC-C/C2. The insert is the spectra of SiOC ceramic.
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Fig. 4. SEM images: inside of PyC-CFt/Sigel (a), PyC-CFt/Sigel after 1050 °C heat treating without PyC (b) and SiOC-C/C preform 2 (c); EDS and element content of spot1 (d), spot2 (e) and spot3 (f) corresponding (a), (b) and (c) respectively. The inserts are images at different magnifications.
testing. The dimensions of the samples were 10 mm × 10 mm × 10 mm for the ILSS test. The Archimedes method was applied to measure the density of the as-prepared samples. The flexural strength by three-pointbend testing was conducted by an electronic universal testing machine (CMT 5304, Suns Co. China), with the across-head speed of 0.5 mm/min and a span of 25 mm at room temperature. And all of the testing, at least five samples were tested.
3. Results and discussion 3.1. Morphology and structural SEM images of the surfaces and cross sections from CFt and PyC-CFt were shown in Fig. 2. By comparing the surfaces of CFt and PyC-CFt, XY axis of the PyC-CFt surface is rougher than CFt, coated with PyC with thickness about 200 nm from the Z axis.
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treating without PyC and SiOC-C/C preform 2 (Fig. 4(a–c)) were captured using SEM. In Fig. 4(a), Si gel grains of different sizes uniformly dispersed in the CFts and attached to the fibers as the result of the dehydration of Si sol. The corresponding EDS spectrum (Fig. 4(d)) showed that the three main ingredients were Si, O and C, whose ratio of atoms number was about 1:2:5. SiOC-C/C preform 2 was prepared by heat-treating PyC-CFt/Sigel at 1050 °C for 2 h, but CH4 was not introduced in this process. As Fig. 4(b) shown, irregular bulk silica materials also attached the fibers, with some irregular round pores on it. The magnified image (top right corner) indicated that the silica materials presented the varying aperture size porous appearance. This morphology was similar to the Si-O-C materials which prepared through a sol-gel method by Liu et al. [30]. From EDS analysis (Fig. 4(e)), these porous materials also contained O, Si and C elements. Compared with PyC-CFt/Sigel, carbon content was decreased. O and Si with an atom ratio of 2:1 were major components. The SiOC ceramic was shown in SEM image of Fig. 4(c), where SiOC phase and fiber phase mixed together and accompanied with many microcracks. From the enlarged graph in Fig. 4(f), the internal was full of rod-like structural, which were composed of elemental Si, O and C revealed by EDS analysis. Due to CVI process, carbon from PyC matrix and SiOC accounted for a major fraction of the materials (76.99% of the total, Atomic). According to the above images, a hypothesis of the possible formation mechanism of the rod-like SiOC ceramic was hereby drawn. As Fig. 5 shown, firstly, with the gradually increasing temperature, Si gel gradually formed a porous (SiO)x ceramic, at the same time, the free carbon from CH4 decomposition during CVI enclosed around the material infiltrated into the microspores. Then, the interaction among C and SiO groups occurred, and eventually formed the amorphousstructured SiOC material. A similar mechanism analysis had also been discussed by Li et al. [31]. Raman spectroscopy is used to evaluate the degree of orderliness or randomness within the carbon materials, because quantitative analysis can be achieved by the intensity ratio of the D peak and G peak in Raman spectra [32]. All the spectra taken from the three composites exhibited two prominent peaks, which were the disorder-induced D mode (~1360 cm−1) and the tangential G mode (~1580 cm−1), as shown in Fig. 6. From the spectra of C/C composites, we can see the maximal intensity ratio of the D peak and G peak, which illustrated that there were many defects in the C/C composites. These defects may be from the interfaces and densification process. Due to the typical vibration modes in carbon based materials [33,34], the Raman spectrum of the SiOC materials [35] also showed the presence of two high intensity bands at about 1360 and 1580 cm−1 which corresponded D peak and G peak respectively. Therefore there was no obvious peak location shift among SiOC-C/C1 (b), SiOC-C/C2 (c) and C/C (a). But the ratio of the intensities of the D and G bands (ID/IG) were determined to be 1.35, 1.17 and 1.09, respectively. Comparing these ratio with C/C, SiOC-C/C1 and SiOC-C/C2, the results suggest that there were a better bond and fewer defects among PyC matrix. As can be seen in curve c in Fig. 6, SiOC-C/C2 had the lowest value of ID/IG, which may be contributed by the PyC interface layer among fibers and carbon matrixes. With the bond density within the interlayer increased, the density increase and the defects decreased [36,37]. In a nutshell, high density and few defects could lead to a significant increase in mechanical strength of composites.
Fig. 5. Schematic structure of SiOC material.
Fig. 6. Raman spectra for composites (a) C/C (b) SiOC-C/C1 (c) SiOC-C/C2.
Fig. 3 showed the XRD patterns of C/C (a), SiOC-C/C1 (b) and SiOCC/C2 (c), and the insert was the spectra of SiOC ceramic. From the patterns, one obvious diffraction peak at about 26° was observed, which was assigned to (002) diffraction plane of carbon fiber [25,26]. The spectra of the SiOC ceramics with a broad peak between 20° and 30° indicated an entirely amorphous structure, which demonstrated the presence of non-crystalline silica [27–29]. The overlapping peak between C and SiOC lead to no obvious difference among the three samples. Although C and SiOC phases coexist in the sample, the structure did not changed muchly. To investigate the growth mechanism of rod-like SiOC ceramic, cross-section images of PyC-CFt/Sigel, PyC-CFt/Sigel after 1050 °C heat
3.2. Physical and mechanical properties Fig. 7 showed the three-point bending test results of C/C, SiOC-C/
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Fig. 7. Three-point bending stress-strain curves (a) average bending strengths (b) and bending maximum loading (c) of the samples.
C1 and SiOC-C/C2. The distinctive behaviors in these samples were presented by the curves with different slopes in Fig. 7(a). Compared with C/C, the bending strengths of SiOC-C/C1 was up slightly, which means that rod-like SiOC ceramic can improve the bending strengths. But without optimizing carbon layer interface, the mechanical properties could not be improved obviously. SiOC-C/C2 showed the highest slope than the rest samples. This observation suggests that both introducing rod-like SiOC ceramic and optimizing interface by PyC can enhance the bending performance. We hereby speculated that fractures could gradually occur during the loading process. Firstly, some fibers fractured and failed, which would led to fine incline of the bending performance. It was obvious that introducing SiOC can increase the brittleness of SiOC-C/C1 and SiOC-C/C2 materials, which then will break under higher loading. In addition, among fiber bundles, SiOC functions as a strengthening agent, which could then fix the fiber bundles and keep them from separating. The bending strengths of SiOCC/C1 and SiOC-C/C2 composites have improved by nearly 37.22% and 118.85%, compared with that of C/C composites, and the bending maximum loading had also increased by 34.60% and 121.42%, respectively.
Fig. 8(a) showed the ILSS stress-strain curves obtained corresponding to C/C, SiOC-C/C1 and SiOC-C/C2 specimens. The curves displayed a clear the specimens' response behaviors. As seen by the difference in slopes, response curves of SiOC-C/C2 specimens showed a higher slope than those of the C/C and SiOC-C/C1 specimens. This observation suggests that the former composites were stiffer than the latter ones. Fig. 8(b) and (c) showed the average shear strengths and shear maximum loading of the samples. Apparently the shear strength (Fig. 8(b)) and shear maximum loading (Fig. 8(c)) of SiOC-C/C1 and SiOC-C/C2 were significantly increased after the in-situ growth of SiOC ceramic in the C/C composites, especially for the composites optimized by the PyC interlayer. Compared with C/C composites, the results indicated that the shear strengths of SiOC-C/C1 and SiOC-C/C2 have improved by 45.7% and 92.2%, respectively. Shear maximum loading of SiOC-C/C1 and SiOC-C/C2 also have improved by 23.44% and 57.46%, respectively. 3.3. Fracture behavior In order to further analyze the enhancement mechanism of C/C
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Fig. 8. ILSS stress-strain curves (a) average shear strengths (b) and shear maximum loading (c) of the samples.
composites by in-situ grown rod-like SiOC ceramic, the fracture surface of the bend tested samples was investigated. The results were shown in Fig. 9, which displayed fracture surface of C/C, and different-magnification SEM images of SiOC-reinforced C/C composites. When the flexural stress is loaded on the composite samples, the strength of the composites is mainly depended on the strength of CFs. Based on the data of Fig. 7 and SEM Fig. 9(a), the fracture process in C/C during bending test can be inferred as follows: when the bending stress was loaded on the C/C samples, destructive cracks originated somewhere within the matrix at the most critical flaws [38]], then propagated along the poor F/M interface, which led to some fibers being pulled out from the PyC. These fibers first were destroyed in failure, which substantially weakened the overall performance of the materials. As the loading increased, the cracks from PyC matrix expanded, and failure rates of fibers continued until thorough destruction disruption [16]. As for SiOC-C/C2, the fracture surface was flat and with nearly absented of fiber pullout (Fig. 9(b)). This fracture surface could be attributed to the strongly-enhanced cohesion in matrix and also the powerful mechanical interlocking at F/M interface, which contributed to a straight and total break rather than step-by-step. Additionally, from the magnification SEM, SiOC ceramic rods pulling out from PyC matrix (Fig. 9(c)) and bridging between fiber and PyC (Fig. 9(d)) could also be found.
Bridging could effectively inhibit cracks to extend. Because inhibited crack expansion and reinforced phase pullout require a great amount of fractured energy consumption during the failure process [15,39], this mechanism in turn increased the flexural strength and drove a gradual release of the flexural stress. The fracture toughness of SiOC-C/C2 also had a greatly improved compared with C/C. The corresponding EDS spectrums (Fig. 9(e) and (f)) from spot 1 (Fig. 9(a)) and spot 2 (Fig. 9(b)) proved the existence of SiOC. When the interlaminar shear stress is loaded on the composite samples, the shear strength is mainly depended on both the matrix cohesion and F/M interfacial bonding strength [38]. Fig. 10 presented the shearing fracture surface of three composites. Firstly, Fig. 10(a–c) showed fracture surfaces paralleled to the fibers. As seen in Fig. 10a, the fractured surface of C/C showed flat and few PyC existed on the surface of fibers (inset), which suggested a poor bonding strength between PyC matrix and Fibers. As for the SiOC-C/C1 (Fig. 10(b)), matrix delaminating from fibers was inhibited and abundant damaged SiOC can be clearly observed in the shearing fracture surface, indicating that the interfacial strength between fibers and PyC was stronger than C/C. This phenomenon also prevented crack expansion along the interface between fiber and PyC matrix. Thus this mechanism led to cracks deflection and fractured energy consumption. Especially for SiOC-C/C2
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Fig. 9. SEM images of the bending failure specimens from C/C (a) and SiOC-C/C2 (different-magnification, b, c and d), EDS spectrums from spot 1 (e) and spot 2 (f).
(Fig. 10(c)), because of optimized the interlayer of F/M, PyC tightly attached on the surface of fibers. Therefore, C/Cs the matrix cohesion is lower than F/M interfacial bonding strength. This observation provided a direct evidence that the carbon interface layer can significantly improve interfacial bonding strength between fibers and PyC matrix. It also could be testified by the Fig. 10(d–f) which showed the fracture
surfaces perpendicular to the fibers of C/C, SiOC-C/C1 and SiOC-C/C2 respectively. Compared with C/C, the adhesiveness of SiOC-C/C1 and SiOC-C/C2 were both improved, and SiOC-C/C2 was the best among them. At the same time, the cracks in the matrix became less obvious, and the defects in this material could be reduced drastically, which also was proved by Raman result.
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Fig. 10. (a) SEM images of the shearing fracture surfaces of the three composites: (a–c) parallel to the fibers and (d–f) perpendicular to the fibers (insets are the magnified images) of the C/C, SiOC-C/C1 and SiOC-C/C2 respectively.
TEM images of the F/M and M/M interfaces were reported in Fig. 11. A layered structure was observed in the SiOC-C/C2 composites, and around the fiber, PyC phase and SiOC particles can also be found. Already at low magnification it could be seen that cracking was induced at the F/M interface, owing to the external force during the sample preparation. But the crack stopped propagating when it encountered the SiOC particle. So it can be inferred that the SiOC can prevent crack growth from the failure occurring along the interfaces under mechanical loading. SiOC can strengthen the interfaces of the SiOC-C/C2 composites, which made the material have stronger mechanical
properties and not easily to be destroyed. Even a lot of cracked SiOC ceramic rods can be observed in the matrix fracture of SiOC-C/C2 (Fig. 12). So introducing rod-like SiOC ceramic into C/C composites could greatly improve both the matrix cohesion and F/M interfacial bonding strength. For SiOC-C/C2, on one hand, an excellent interface bonding can effectively inhibit the initiation and propagation of cracks along the F/M interface. On the other hand, SiOC ceramic rods' own strength and high matrix cohesion all consume the stress. It also explained why the increment in ILSS of SiOCC/C2 (that is 92.2%) was higher than that of SiOC-C/C1 (that is 45.7%).
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Fig. 11. Typical TEM micrographs of the F/M and M/M interfaces domain in the SiOC-C/C2 composites: (a) low magnification and (b) high magnification with the EDS spectrum.
Fig. 12. SEM images cross-section of failure specimens of SiOC-C/C2.
4. Conclusions
mechanical strength to replace traditional C/C in industries.
Rod-like SiOC ceramics were in-situ grown in the C/C composites by CVI reaction method using the porous properties of (SiO)x ceramic. In order to optimize F/M interface, strength and protect fibers, about 200 nm carbon layer has been deposited on the CFs surface at the beginning. Three-point bending and ILSS testing results showed that the above-mentioned material named SiOC-C/C2 have nearly 118.85% and 92.2% increase improvements compare to C/C composites in bending and shear strengths respectively. SEM morphologies of fracture surfaces of failure composites revealed that the synergistic effects of grown rodlike SiOC ceramic and carbon layer interface along with optimized F/M interfaces not only efficiently improved the matrix cohesion, but also the F/M interfacial bonding strength, which significantly decrease cracks generated and inhibited the crack propagation. This work opens up a possibility to produce SiOC-reinforced C/Cs with excellent
Author statement I have revised the manuscript of “Optimizing PyC matrix interface to improve mechanical properties of carbon/carbon composites by rodlike SiOC ceramic”. I have approved the final version to be published. All authors agree to publish this article, and all the dates are original. I certify that I did not submit it to other journals. Declaration of competing interest We declare that we have no financial and personal relationships with other people or organizations that can inappropriately influence
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our work. There is no professional or other personal interest of any nature or kind in any product, service or company that could be construed as influencing the position presented in, or the review of, the manuscript entitled” Optimizing PyC matrix interface to improve mechanical properties of carbon/carbon composites by rod-like SiOC ceramic”.
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