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C composites
Journal of Alloys and Compounds 726 (2017) 866e874
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Journal of Alloys and Compounds journal homepage: http://www.elsevier.com/locate/jalcom
Effect of in-situ grown SiC nanowires on the mechanical properties of HfC-ZrB2-SiC modified C/C composites Weiyan Wang, Qiangang Fu*, Biyi Tan State Key Laboratory of Solidification Processing, Carbon/Carbon Composites Research Center, Northwestern Polytechnical University, Xi'an 710072, PR China
a r t i c l e i n f o
a b s t r a c t
Article history: Received 31 May 2017 Received in revised form 7 August 2017 Accepted 8 August 2017 Available online 9 August 2017
In order to improve the strength and toughness of the C/C-HfC-ZrB2-SiC composites, uniformly distributed SiC nanowires were in situ grown in C/C preform using a simple precursor infiltration pyrolysis method with ferrocene as catalyst. The SiC nanowires could greatly refine the size of the ultrahigh temperature ceramic (HfC-ZrB2-SiC) particles, and make the ceramic particles interconnected with the unitary body. After introducing SiC nanowires, the flexural strength, flexural modulus and fracture toughness of the composites increased by 21%, 29% and 28% respectively, due to interfacial weakening, the deflection of cracks induced by SiC nanowires and the pullout of SiC nanowires. The C/CHfC-ZrB2-SiC composites with SiC nanowires also displayed good erosion resistance under high velocity Al2O3 particle erosion. © 2017 Elsevier B.V. All rights reserved.
Keywords: Modified C/C composites Precursor infiltration pyrolysis SiC nanowires Ultra-high temperature ceramic Mechanical properties
1. Introduction Carbon fiber reinforced carbon matrix (C/C) composites have a great potential for high performance structural materials due to their low density, low coefficient of thermal expansion, high specific strength/modulus, excellent thermal shock resistance and good mechanical properties at elevated temperature [1e3]. These outstanding properties make them the leading structural materials used in high temperature environment such as rocket engines, nose tips, leading edges and other thermal protection systems for space [4,5]. However, C/C composites are vulnerable when ablated at ultrahigh temperature (over 2000 C) and under high velocity particle erosion, which restricts their wide applications in aerospace vehicles [6,7]. Numerous attempts have been made on addressing this issue, and it is proved that introducing ultrahigh temperature ceramics (UHTCs) into C/C composites can play a positive role on the ablation resistance improvement of the composites [8e10]. ZrB2 as one of the UHTCs has high melting point (3313 K), high hardness, low volatilization rate, good chemical stability, and good ablation resistance [11]. Unfortunately, the oxidation resistance of ZrB2 can
* Corresponding author. E-mail address: [email protected] (Q. Fu). http://dx.doi.org/10.1016/j.jallcom.2017.08.060 0925-8388/© 2017 Elsevier B.V. All rights reserved.
be affected by temperature easily. ZrB2 has poor oxidation resistance above 1373 K [12]. The mixture of SiC and ZrB2 could form a silicate glass after oxidation between 1273 K and 2073 K, which could increase the oxidation resistance of ZrB2 at elevated temperature. Besides, SiC is usually used as a transition to decrease the mismatch of CTE between carbon and other UHTCs because of its lower coefficient of thermal expansion (CTE) [13e15]. Up to now, C/ C-ZrB2-SiC composites with elevated ablation property have been successfully synthesized [16]. Hafnium carbide (HfC) has the highest melting point (4163 K) among the single compounds and exhibits good ablation resistance, low evaporation (vapor) pressure, and good chemical inertness. In addition, the oxide of HfC (HfO2) has an appropriately high melting point and a relatively lower vapor pressure. Therefore, HfC is a potential candidate to modify C/C composites against ablation [17,18]. Nevertheless, the low fracture toughness of the ceramics and the mismatch of the CTE between C/ C and the UHTCs may lead to the formation of penetrating cracks and end up showing a brittle behavior and poor thermal shock resistance [19]. One-dimensional nanostructures have great potential as reinforcements for composite materials because of their significantly greater strength than their bulk counterparts [20e22]. The reported elastic modulus and ultimate bending strength of SiC nanowires are 610e660 GPa and 53.4 GPa, respectively [23]. In the last few years, SiC nanowires have been widely studied as
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reinforcements [24]. Chu et al. [25] introduced the SiC nanoribbon into SiC-Si ceramic coating by pack cementation, and the results show that the fracture toughness of the coating and the interfacial bonding strength of the coating increased by 98% and 99%, respectively. Dong et al. [26] found that the hardness and tensile properties of SiC nanowires/6061Al composite significantly increased by 40% after peak-aging treatment when high content SiC nanowires were added to the 6061Al matrix. However, up to now, little work has been reported on the synthesis and properties of insitu forming SiC nanowires in C/CeUHTC composites. Precursor infiltration pyrolysis (PIP) is an effective method to synthesize ceramics in the porous preform due to the ceramic distribution uniformity and the ceramic component designability [27]. In our previous work [21], SiC nanowires were successfully insitu grown on the carbon fiber by PIP method, which distributed inside the C/C performs uniformly. It is worth attempting to use the in-situ grown SiC nanowires to improve the strength and toughness of the C/CeUHTC composites. In this work, the in-situ grown SiC nanowires were introduced into C/C-UHTC composites by PIP method, using the mixture solution of polycarbosilane (PCS), organic HfC polymer (PHC), organic ZrB2 polymer (PZC) as precursor, and ferrocene as catalyst. The effects of in-situ grown SiC nanowires on the microstructure of the ceramic matrix, mechanical property, fracture toughness and particle erosion of the SiC nanowires (NWs) -reinforced C/C -HfC-ZrB2SiC (HZS) were investigated.
2. Experimental procedure 2.1. Materials preparation 2.5-D needle punched carbon fiber fabric (0.45 g/cm3) was densified through thermal gradient chemical vapor infiltration (TCVI) process to produce porous C/C composites with a density of 1.3 g/cm3. Polycarbosilane (PCS), organic HfC precursor (PHC) and ferrocene ((Fe(C5H5)2), purity: 99.0%) were mixed with a weight
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ratio of 1:2:0.01, and then were dispersed in dimethylbenzene to form a homogenous mixed precursor by ultrasonic dispersion at room temperature for 10 h. Fig. 1 shows the schematic preparation procedures of the SiCNWs-C/C composites by PIP method. Firstly, the mix precursor was infiltrated into the porous C/C composites, and then the C/C preforms were dried in a drying oven for 24 h. Finally, they were undergone a heat treatment at 1400 C for 2 h at an argon atmosphere. SiC nanowires were expected to grow in-situ in the C/C preform during the heat treatment. The as received SiCNWsC/C preforms were subjected to densification process by PIP method. The PHC, PCS and organic ZrB2 precursor (POZ) were mixed with a weight ratio of 2:1:1 to form a mixed precursor. After PIP treatment for 10 times, the SiC nanowire-modified C/C-HfCZrB2-SiC composites were fabricated, marked as NW-C/C-HZS composites. Meanwhile, pure C/C preforms without in-situ grown SiC nanowires were subjected to the same PIP treatment under the same conditions for comparisons, marked as C/C-HZS composites.
2.2. Physical and mechanical property tests Density and open porosity of the specimens were measured by Archimedes principle. Each data point was an average of three values. Three-point-bend tests and compression tests were carried out on an electronic universal testing machine (CMT 5304, Suns Co. China) with a loading rate of 0.5 mm/min at room temperature. The sizes of samples used for above tests were 55 mm 10 mm 4 mm. The final results were achieved by computing the average values of at least three specimens. The flexural strength (sf) and flexural modulus (Ef) were calculated from the data recorded during the tests. In addition, the measurements of the fracture toughness (KIC) of as-prepared composites were obtained through three point single edge notched beam (SENB) testing. The initial crack width and length were 200 mm and 2 mm, respectively, which was obtained using blade cutter manually. The particle erosion test was carried out on the shot blasting machine (1010FK, Jichuan Co, China) using Al2O3 particles with a
Fig. 1. Schematic preparation procedures of SiCNWseC/C composites.
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Fig. 2. The XRD patterns of the two modified C/C composites.
diameter of 1~2 mm. The particle flux is 0.24 m3/min. The size of the samples was 430 mm 10 mm. The inner diameter of the gun tip was 5 mm and the distance between the gun tip and the samples was 20 mm. Linear and mass erosion rates were calculated by the change of thickness and weight before and after erosion test. The ultimate result was the average value of three samples. 2.3. Microstructure characterization The morphology and elements of the samples were analyzed using a field emission scanning electron microscopy (FESEM, ZEISSSUPRA 55, Germany), equipped with energy dispersive X-ray spectroscopy (EDS). The phase compositions of the samples were analyzed by X-ray diffraction (XRD, X'Pert PRO MPD) using Cu as anode material. The microstructures of the as-prepared SiC nanowires was characterized by a transmission electron microscope (TEM), high-resolution transmission electron microscopy (HRTEM) and selected area electron diffraction (SAED) images (FEI TecnaiF30G2, USA). Fourier transform infrared (FT-IR) and Raman shift spectrum at room temperature were recorded by a FT-IR spectrometer (Bruker Tensor-27, Germany), a Raman spectrometer (RMS; Renishaw, UK), respectively. 3. Results and discussion 3.1. Microstructure
diffraction peaks. C, SiC, HfC and ZrB2 exist in both samples, as well as HfO2 and SiO2. The existence of HfO2 and SiO2 could be the result of the incomplete pyrolyzation of the precursor under high temperature heat treatment process. Moreover, the diffraction peaks of NW-C/C-HZS is sharper than those of C/C-HZS, indicating that the addition of SiC nanowires could effectively improve the crystallinity of the ceramics. The distribution and morphology of SiC nanowires in C/C performs will directly affect the subsequent ceramic densification process and the properties of the obtained composites. Fig. 3a shows the SEM image of nanowires inside C/C performs. The distribution of nanowires is uniform inside the preform, attributing to the good infiltration ability of the mixed liquid precursor in porous C/C preform. The nanowires are generally several tens to more than one hundred micrometers in length and randomly oriented with straight morphologies, as confirmed by the higher-magnification SEM image in Fig. 3b. To confirm the structure of the nanowires, FT-IR measurement and Raman scattering were employed to analyze the samples. The typical FT-IR spectrum and Raman shift spectrum are shown in Fig. 4. The absorption peaks at around 810, 1080 and 1610 cm 1 were detected by FT-IR. The strong absorption peak at around 810 cm 1 is stretching vibration of Si-C, indicating that the assynthesized nanowires should be SiC. The other two weak absorption bands centered at 1080 and 1610 cm 1 are attributed to the Si-O and H-O-H bending vibration, which might be caused by the incomplete conversion of the precursors and the adsorbed water. Similarly, the Raman shift spectrum shows that the sharp peaks at 790 cm 1 correspond to the modes of transverse optical (TO) phonons of the 3C-SiC [28] in Fig. 4b. In addition, a down shift appears compared to the previous result [29] (TO mode at 796 cm 1 obtained from bulk 3CeSiC). This should be attributed to the confinement effect and inner stress from the structural defect of achieved nanowires. The low-resolution TEM image of the nanowire is shown in Fig. 5a, which shows that the SiC nanowire is linear with uniform structure of about 150 nm in diameter. The corresponding EDS patterns inserted in Fig. 5a shows that the nanowire is composed of Si and C. From selected area electron diffraction (SAED) pattern (inserted in Fig. 5b) and high-resolution TEM (HRTEM) image (Fig. 5b) of a single SiC nanowire, it can be seen that the as-received SiC nanowire has perfect single crystalline structure, possessing high-density stacking faults in the crystal plane normal to the axis of the nanowire. The formation of these stacking faults could consume lower energy compared with the perfect crystal planes [30]. The nanowire crystal lattice fringe spacing is 0.25 nm, which is the same as the distance between {001} planes in 3C-SiC crystalline.
Fig. 2 shows the XRD patterns of the samples. It can be seen that two kinds of modified C/C composites as prepared have similar
Fig. 3. SEM images of nanowires inside C/C performs.
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Fig. 4. FT-IR absorbance of the as-synthesized nanowires (a) and Raman shift spectrum (b).
Fig. 5. TEM image of the single SiC nanowire with the insert showing the corresponding EDS pattern (a) and HRTEM image of the single SiC nanowire and its corresponding SAED pattern (b).
3.2. Mechanical properties Table 1 shows the density, porosity and mechanical properties of the composites. The value of density and porosity are the average value of three specimens measured by Archimedes principle. The C/ C-HZS composites have an average density of 2.27 ± 0.05 g/cm3, with an average porosity of 12.3 ± 0.4%. The NW-C/C-HZS composites have similar density (2.23 ± 0.06 g/cm3) and lower porosity (11.4 ± 0.3%). The obtained flexural strength and fracture toughness are also summarized in Table 1. The average flexural and fracture toughness of the NW-C/C-HZS composites are up to 172 ± 8 MPa and 9.7 ± 0.4 MPa$m1/2, respectively. Compared to the C/C-HZS composites, the flexural and fracture toughness are enhanced by 21% and 28%, respectively. The results show that the incorporation of SiC nanowires greatly improves the flexural strength and fracture toughness of C/C composites modified by HfC-ZrB2-SiC ceramics. Fig. 6 shows the flexural stress-strain curves of the composites. It is obvious that the curves of two composites exhibit a similar trend. It can be seen that in the initial stage of loading load, both materials have a quasi-elastic strain region, and then the flexural strength rapidly declines after the strain reached the maximum
value. It is noteworthy that the curve of the NW-C/C-HZS composites appears a shock region when the strain reaches the maximum value, which is the result of the alternate effects of interfacial debonding and carbon fiber bridging in ceramic matrix, while the C/C-HZS composites break immediately when strain reach the maximum value. In addition, the curve of the C/C-HZS composites shows some small inflection points during the quasiplastic deformation period, which indicates that microcracks appeared in the ceramic matrix, while the curve of the NW-C/CHZS composites is smoother under the same strain. As shown in Fig. 7, the fracture surfaces of the fabricated composites after three point bending test were observed by SEM. In Fig. 7a, the fracture surface is rather planar and few fibers are pulled out, with obvious brittle fracture characteristics, indicating the poor toughness of the C/C-HZS composites. Fig. 7b shows that the carbon fiber bundles are pulled out from the NW-C/C-HZS composites. There are a large number of SiC nanowires and ceramic particles on the surface of fiber bundles, which indicates that the interface between carbon fiber bundles and ceramic matrix was debonding during the process of carbon fiber bundle pull-out and a large amount of ceramic particles remain on the fiber bundle
Table 1 Density, porosity and mechanical properties of composites.
C/C-HZS NW-C/C-HZS
Density (g/cm3)
Porosity (%)
Flexural Strength (MPa)
Fracture Toughness (MPa$m1/2)
Flexural Modulus (GPa)
2.27 ± 0.05 2.23 ± 0.06
12.3 ± 0.4 11.4 ± 0.3
143 ± 9 172 ± 8
7.6 ± 0.3 9.7 ± 0.4
14.2 ± 0.4 18.3 ± 0.3
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Fig. 6. The flexural stress-strain curves of C/C-HZS (a) and NWeC/C-HZS (b) composites.
surface due to the connection of SiC nanowires. Apparently, the fracture surfaces of these two kinds of modified C/C composites are different. The interfacial analysis of the carbon fiber and ceramic matrix was carried out to further understand the cause of fracture and the effect of SiC nanowires on the modified composites. Carbon fiber bundles of the C/C-HZS and NW-C/C-HZS composites in the fracture surface are shown in Fig. 8a and b, respectively. From Fig. 8a, it can be seen clearly that the carbon fiber is closely bonded with the matrix ceramic, and the matrix consists of bulk ceramic particles. A lot of obvious cracks appear in the ceramic matrix, which is due to decomposition and volume
shrinkage of the precursor during the precursor pyrolysis process. From Fig. 8b, the ceramic particles of the NW-C/C-HZS composites are small and uniformly distributed inside the network of SiC nanowires, and the ceramic particles are interconnected by SiC nanowires so that the matrix becomes a unitary body. Compared with C/C-HZS composites, there are a lot of micro holes in the ceramic matrix instead of obvious cracks, due to the existence of SiC nanowires. This kind of structure might cause a weaker interface between the carbon fiber and ceramic matrix. During the fracture process of composite materials, when it is subjected to a certain loading, some cracks will inevitably be formed in the composites, and then the cracks will gradually extend till they cut through the specimen, finally the composite will completely be broken. Therefore, the fracture process of modified C/C composites is actually the process of crack propagation inside the material. There are two general paths for crack propagation: perpendicular to the fiber direction and parallel to the fiber direction. For C/C-HZS composites, due to the strong interfacial bonding between carbon fiber bundles and ceramic matrix, the cracks can not be deflected when the load is transferred to the ceramicpyrolytic carbon interface, and the stress is concentrated at the interface. When the stress gradually increases to the maximum value, the cracks expand along the interface (parallel to the fiber direction). Perpendicular to the fiber direction, the cracks pass directly through the pyrolytic carbon to the fiber, which leads to brittle fracture of pyrolytic carbon matrix and fibers (Fig. 9a). For NW-C/C-HZS composites, the interfacial bonding between the ceramic matrix and the pyrolytic carbon is slightly weakened due to the introduction of SiC nanowires. When the crack propagates to the fiber, the interface starts to debond to prevent the crack from propagating forward. The crack is forced to change its original prolongation direction and extend along the interface by the effect of stress concentration, which eventually causes the debonding of the ceramic matrix and pyrolytic carbon, increasing the fracture
Fig. 7. The fracture surfaces of the fabricated composites after three point bending test. (a) C/C-HZS composites; (b) NW-C/C-HZS composites.
Fig. 8. The interface of the carbon fiber and ceramic matrix of the composites in the fracture surface of C/C-HZS composites (a) and NW-C/C-HZS composites (b).
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Fig. 9. The cracks propagation of the C/C-HZS composites (a) and NW-C/C-HZS composites (b); the cracks propagation in the ceramic matrix of the C/C-HZS composites (c) and the NW-C/C-HZS composites (d).
stress. From Fig. 9b, it is observed that the pyrolytic carbon layer appears delamination phenomenon and the stepped layer is pulled out. Due to the debonding of ceramic matrix and pyrolytic carbon layer, the stress could be released partially. As the cracks encounter obstacles again, the stress is concentrated at the interface between the carbon fiber and the pyrolytic carbon layer since carbon fiber could still bear the load. The mismatch of the coefficient of thermal expansion between the carbon fiber and the pyrolytic carbon layer will result in the generation of circumferential microcracks. The propagation path of circumferential microcracks might be changed along the interface direction when the crack tips are distributed vertically at circumferential microcracks, resulting in the debonding of the interface between the carbon fiber and the pyrolytic carbon layer and the delamination of the pyrolytic carbon layer under the effect of stress. Finally, the carbon fibers break, and they are pulled out from the pyrolytic carbon layer. From Fig. 9c, the ceramic matrix breaks directly because of the low fracture toughness of the ceramics. While the path of the crack propagation is deflected due to the obstacle of the SiC nanowires' network system and ceramic particles' refinement (Fig. 9d). In the ceramic matrix of
the NW-C/C-HZS composites, the ceramic particles are refined because of the introduction of SiC nanowires, and ceramic particles are uniformly distributed in the network system of SiC nanowires and have a more solid connection. When the cracks propagate to the ceramic matrix, the refinement of the ceramic particles can product more obstacles for the crack propagation process, resulting in the deflection of the path of cracks propagation, and the fracture of the SiC nanowires and the deflection of the crack will dissipate extra energy. The effect of SiC nanowires in the fracture process of the NW-C/ C-HZS composites was further investigated. As shown in Fig. 10a, it is clearly observed that there are a large number of SiC nanowires in the interstices of the ceramic particles, which are connected to each other, and ceramic particles are also connected by SiC nanowires. Under the connection of the SiC nanowires, the ceramic matrix becomes an interconnected whole, which enhances the strength of the matrix, thus the mechanical properties of the composites are improved as well. In addition, SiC nanowires formed in situ in the NW-C/C-HZS composites appear in the ceramic blocks after the multiple densification process, which improves the toughness of
Fig. 10. A representative image of ceramic particles bridging by the nanowire (a) numerous pullout or debonding nanowires (b).
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Table 2 Density, porosity and particle erosion performance of composites.
C/C-HZS NW-C/C-HZS
Density (g/cm3)
Porosity (%)
Linear erosion rate (10
2.41 ± 0.05 2.35 ± 0.04
11.6 ± 0.8 12.5 ± 0.6
28 ± 5 11 ± 3
3
mm/s)
Mass erosion rate (10
4
g/s)
3.6 ± 0.6 2.3 ± 0.4
Fig. 11. SEM images of the surface of composites after particle erosion. (a) C/C-HZS composites; (b) NW-C/C-HZS composites; (c) the holes of ceramic matrix of the C/C-HZS composites; (d) the craters of ceramic matrix of the NW-C/C-HZS composites.
ceramic. From Fig. 10b, the pulling out of SiC nanowires can be observed on the fracture surface of ceramic blocks. The stress increases and the microcracks expand when the matrix is under load. Due to the existence of SiC nanowires in some of the ceramic blocks, the propagation of cracks through the ceramic are impeded. The SiC nanowires debond from the ceramic and then pull out, which consumes much energy, so that the ceramic blocks will not crack directly. Therefore, SiC nanowires have a toughening effect on the matrix in this process.
3.3. The particle erosion performance of the composites Table 2 shows the density and the porosity of the two composites after subjected to Al2O3 particle erosion. The C/C-HZS composites have an average density of 2.41 ± 0.05 g/cm3, which corresponds to a porosity of 11.6 ± 0.8%. The composites with SiC nanowires have similar density (2.35 ± 0.04 g/cm3) and similar porosity (12.5 ± 0.6%) with those without SiC nanowires. It shows that the NW-C/C-HZS composites have lower linear and mass erosion rates than the C/C-HZS composites. It is meaningful to investigate the erosion surfaces of different composites to understand the effect of SiC nanowires on the particle erosion performance. Fig. 11 shows the morphology of two modified C/C composites after particle erosion. As shown in Fig. 11a, the surface of the C/C-HZS composites not only has many craters, but also leaves a lot of holes. By further magnification to observe the holes (Fig. 11c), it is clear to see that the Al2O3 particles directly puncture the short-cut fiber web of the ceramic matrix, and the next non-woven layer can be seen through the hole. Comparatively,
the surface of the NW-C/C-HZS composites also leaves many craters, but just few pore after particle erosion (Fig. 11b). Fig. 11d shows the magnification SEM image of the crater of the NW-C/CHZS composite. Compared to the craters of the C/C-HZS composites, the size of the craters of the NW-C/C-HZS composites is smaller, and the short-cut fiber web is not punctured by high energy Al2O3 particles. To further investigate the erosion morphologies of two modified composites, the craters were observed by SEM. Fig. 12 shows the craters on the ceramic matrix of two modified C/C composites after particle erosion. As shown in Fig. 12a, the big bulk ceramic particles in the ceramic matrix of the C/C-HZS composite are crushed and peeled off. From Fig. 12b, it is clear that the size of ceramic particles of the NW-C/C-HZS composite is much smaller than that of the C/CHZS composite. From Fig. 12c and d, the size of ceramic particles of the NW-C/C-HZS composites is about 1e4 mm and three kinds of ceramic particles are uniformly distributed, embedded in each other. In addition, SiC nanowires can also be found between the ceramic particles in Fig. 12d, which connect the ceramic particles into a whole. During the particle erosion process, ceramic particles can effectively transfer energy to each other through SiC nanowires, so that the ceramic matrix has better erosion resistance.
4. Conclusions The SiC nanowires reinforced C/C-HfC-ZrB2-SiC composites were fabricated by a multistep technique of precursor infiltration and pyrolysis process. Due to the introduction of SiC nanowires, the size of the ultra-high temperature ceramic particles were refined
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Fig. 12. The craters leave on the ceramic matrix after particle erosion of C/C-HZS composites (a) and NW-C/C-HZS composites (b); (c) and (d): high resolution images of (b).
greatly, and the ceramic matrix becomes a unitary body with the interconnection of SiC nanowires, hence the NW-C/C-HZS composites displayed good particle erosion resistance. The flexural strength, flexural modulus and fracture toughness of the NW-C/CHZS composites increased by 21%, 29% and 28%, respectively compared to the C/C-HZS composites without SiC nanowires. The improvement of toughness and strength of the HZS ceramics is attributed to the pullout of carbon fiber, the deflection of cracks induced by nanowires and the pullout of SiC nanowires. Acknowledgments This work was supported by National Natural Science Foundation of China [grant numbers 51521061, 51432008 and 51572223] and the “111” Project [grant number B08040]. References [1] X. Shen, K. Li, H. Li, et al., Microstructure and ablation properties of zirconium carbide doped carbon/carbon composites, Carbon 48 (2) (2010) 344e351. [2] L. Gaab, D. Koch, G. Grathwohl, Effects of thermal and thermomechanical induced mechanical changes of C/C composites, Carbon 48 (10) (2010) 2980e2988. [3] J.C. Rietsch, J. Dentzer, A. Dufour, F. Schnell, L. Vidal, P. Jacque, et al., Characterizations of C/C composites and wear debris after heavy braking demands, Carbon 47 (1) (2009) 85e93. [4] Y.L. Zhang, Z. Hu, B. Yang, et al., Effect of pre-oxidation on the ablation resistance of ZrB 2 eSiC coating for SiC-coated carbon/carbon composites, Ceram. Int. 41 (2) (2015) 2582e2589. [5 ] T.L. Dhami, O.P. Bahl, B.R. Awasthy, Oxidation-resistant carbon-carbon composites up to 1700 C, Carbon 33 (4) (1995) 479e490. [6] Q.G. Fu, L. Zhuang, Q.W. Ren, et al., Carbon nanotube-toughened interlocking buffer layer to improve the adhesion strength and thermal shock resistance of SiC coating for C/CeZrCeSiC composites, J. Materiomics 1 (3) (2015) 245e252. [7] D.D. Jayaseelan, R.G.D. S a, P. Brown, et al., Reactive infiltration processing (RIP) of ultra high temperature ceramics (UHTC) into porous C/C composite tubes, J. Eur. Ceram. Soc. 31 (3) (2011) 361e368. [8] L. Liu, H. Li, W. Feng, et al., Effect of surface ablation products on the ablation resistance of C/CeSiC composites under oxyacetylene torch, Corros. Sci. 67 (2013) 60e66. [9] K.Z. Li, X. Jing, Q.G. Fu, et al., Effects of porous C/C density on the densification behavior and ablation property of C/CeZrCeSiC composites, Carbon 57 (3) (2013) 161e168.
[10] L. Zhuang, Q.G. Fu, B.Y. Tan, et al., Ablation behaviour of C/C and C/CeZrCeSiC composites with cone-shaped holes under an oxyacetylene flame, Corros. Sci. 102 (2016) 84e92. [11] S.R. Levine, E.J. Opila, M.C. Halbig, et al., Evaluation of ultra-high temperature ceramics foraeropropulsion use, J. Eur. Ceram. Soc. 22 (14e15) (2002) 2757e2767. [12] E. Padovano, C. Badini, E. Celasco, et al., Oxidation behavior of ZrB 2/SiC laminates: effect of composition on microstructure and mechanical strength, J. Eur. Ceram. Soc. 35 (6) (2015) 1699e1714. [13] W.G. Fahrenholtz, G.E. Hilmas, I.G. Talmy, et al., Refractory diborides of zirconium and hafnium, J. Am. Ceram. Soc. 90 (5) (2007) 1347e1364. [14] E. Jung, J. Kim, S. Jung, et al., ChemInform abstract: synthesis of ZrB2 powders by carbothermal and borothermal reduction, J. Alloys Compd. 538 (42) (2012) 164e168. [15] J. Ren, Y. Zhang, P. Zhang, et al., UHTC coating reinforced by HfC nanowires against ablation for C/C composites 311 (2017) 191e198. [16] H.J. Li, X.Y. Yao, Y.L. Zhang, et al., Effect of heat flux on ablation behaviour and mechanism of C/CeZrB 2 eSiC composite under oxyacetylene torch flame, Corros. Sci. 74 (4) (2013) 265e270. [17] L. Xue, Z.A. Su, X. Yang, et al., Microstructure and ablation behavior of C/CeHfC composites prepared by precursor infiltration and pyrolysis, Corros. Sci. 94 (2015) 165e170. [18] W. Tan, K. Li, H. Li, et al., Ablation behavior and mechanism of C/C-HfC-SiC composites, Vacuum 116 (9) (2015) 124e129. [19] L. Li, H. Li, Y. Li, et al., A SiC-ZrB 2 -ZrC coating toughened by electrophoretically-deposited SiC nanowires to protect C/C composites against thermal shock and oxidation, Appl. Surf. Sci. 349 (2015) 465e471. [20] Y. Jia, Z. Chen, J. Jiang, et al., Free-standing all-nanoparticle thin fibers: a novel nanostructure bridging zero-and one-dimensional nanoscale features, Adv. Mater. 21 (3) (2009) 314e319. [21] Y. Wang, H. Wu, Microstructure of friction surface developed on carbon fibre reinforced carbonesilicon carbide (C f/CeSiC), J. Eur. Ceram. Soc. 32 (12) (2012) 3509e3519. [22] Q.G. Fu, B.Y. Tan, L. Zhuang, et al., Significant improvement of mechanical properties of carbon/carbon composites by in situ growth of SiC nanowires, Mater. Sci. Eng. A 672 (2016) 121e128. [23] E.W. Wong, P.E. Sheehan, C.M. Lieber, Nanobeam mechanics: elasticity, strength, and toughness of nanorods and nanotubes, Science 277 (5334) (1997) 1971e1975. [24] W. Yang, H. Araki, A. Kohyama, et al., Process and mechanical properties of in situ silicon carbide-nanowire-reinforced chemical vapor infiltrated silicon carbide/silicon carbide composite, J. Am. Ceram. Soc. 87 (87) (2010) 1720e1725. [25] Y. Chu, H. Li, H. Luo, et al., Oxidation protection of carbon/carbon composites by a novel SiC nanoribbon-reinforced SiCeSi ceramic coating, Corros. Sci. 92 (2014) 272e279. [26] R. Dong, W. Yang, P. Wu, et al., High content SiC nanowires reinforced Al composite with high strength and plasticity, Mater. Sci. Eng. A 630 (2015) 8e12.
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[27] B. Pei, Y. Zhu, M. Yuan, et al., Effect of in situ grown SiC nanowires on microstructure and mechanical properties of C/SiC composites, Ceram. Int. 40 (4) (2014) 5191e5195. [28] M. Bechelany, A. Brioude, D. Cornu, et al., A Raman spectroscopy study of individual SiC nanowires, Adv. Funct. Mater. 17 (6) (2007) 939e943.
[29] D. Olego, M. Cardona, Pressure dependence of Raman phonons of Ge and 3CSiC, Phys. Rev. B Condens. Matter 25 (2) (1982) 1151e1160. [30] W.M. Vetter, M. Dudley, Characterization of defects in 3C-silicon carbide crystals, J. Cryst. Growth 260 (1) (2004) 201e208.
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Journal of Alloys and Compounds journal homepage: http://www.elsevier.com/locate/jalcom
Effect of in-situ grown SiC nanowires on the mechanical properties of HfC-ZrB2-SiC modified C/C composites Weiyan Wang, Qiangang Fu*, Biyi Tan State Key Laboratory of Solidification Processing, Carbon/Carbon Composites Research Center, Northwestern Polytechnical University, Xi'an 710072, PR China
a r t i c l e i n f o
a b s t r a c t
Article history: Received 31 May 2017 Received in revised form 7 August 2017 Accepted 8 August 2017 Available online 9 August 2017
In order to improve the strength and toughness of the C/C-HfC-ZrB2-SiC composites, uniformly distributed SiC nanowires were in situ grown in C/C preform using a simple precursor infiltration pyrolysis method with ferrocene as catalyst. The SiC nanowires could greatly refine the size of the ultrahigh temperature ceramic (HfC-ZrB2-SiC) particles, and make the ceramic particles interconnected with the unitary body. After introducing SiC nanowires, the flexural strength, flexural modulus and fracture toughness of the composites increased by 21%, 29% and 28% respectively, due to interfacial weakening, the deflection of cracks induced by SiC nanowires and the pullout of SiC nanowires. The C/CHfC-ZrB2-SiC composites with SiC nanowires also displayed good erosion resistance under high velocity Al2O3 particle erosion. © 2017 Elsevier B.V. All rights reserved.
Keywords: Modified C/C composites Precursor infiltration pyrolysis SiC nanowires Ultra-high temperature ceramic Mechanical properties
1. Introduction Carbon fiber reinforced carbon matrix (C/C) composites have a great potential for high performance structural materials due to their low density, low coefficient of thermal expansion, high specific strength/modulus, excellent thermal shock resistance and good mechanical properties at elevated temperature [1e3]. These outstanding properties make them the leading structural materials used in high temperature environment such as rocket engines, nose tips, leading edges and other thermal protection systems for space [4,5]. However, C/C composites are vulnerable when ablated at ultrahigh temperature (over 2000 C) and under high velocity particle erosion, which restricts their wide applications in aerospace vehicles [6,7]. Numerous attempts have been made on addressing this issue, and it is proved that introducing ultrahigh temperature ceramics (UHTCs) into C/C composites can play a positive role on the ablation resistance improvement of the composites [8e10]. ZrB2 as one of the UHTCs has high melting point (3313 K), high hardness, low volatilization rate, good chemical stability, and good ablation resistance [11]. Unfortunately, the oxidation resistance of ZrB2 can
* Corresponding author. E-mail address: [email protected] (Q. Fu). http://dx.doi.org/10.1016/j.jallcom.2017.08.060 0925-8388/© 2017 Elsevier B.V. All rights reserved.
be affected by temperature easily. ZrB2 has poor oxidation resistance above 1373 K [12]. The mixture of SiC and ZrB2 could form a silicate glass after oxidation between 1273 K and 2073 K, which could increase the oxidation resistance of ZrB2 at elevated temperature. Besides, SiC is usually used as a transition to decrease the mismatch of CTE between carbon and other UHTCs because of its lower coefficient of thermal expansion (CTE) [13e15]. Up to now, C/ C-ZrB2-SiC composites with elevated ablation property have been successfully synthesized [16]. Hafnium carbide (HfC) has the highest melting point (4163 K) among the single compounds and exhibits good ablation resistance, low evaporation (vapor) pressure, and good chemical inertness. In addition, the oxide of HfC (HfO2) has an appropriately high melting point and a relatively lower vapor pressure. Therefore, HfC is a potential candidate to modify C/C composites against ablation [17,18]. Nevertheless, the low fracture toughness of the ceramics and the mismatch of the CTE between C/ C and the UHTCs may lead to the formation of penetrating cracks and end up showing a brittle behavior and poor thermal shock resistance [19]. One-dimensional nanostructures have great potential as reinforcements for composite materials because of their significantly greater strength than their bulk counterparts [20e22]. The reported elastic modulus and ultimate bending strength of SiC nanowires are 610e660 GPa and 53.4 GPa, respectively [23]. In the last few years, SiC nanowires have been widely studied as
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reinforcements [24]. Chu et al. [25] introduced the SiC nanoribbon into SiC-Si ceramic coating by pack cementation, and the results show that the fracture toughness of the coating and the interfacial bonding strength of the coating increased by 98% and 99%, respectively. Dong et al. [26] found that the hardness and tensile properties of SiC nanowires/6061Al composite significantly increased by 40% after peak-aging treatment when high content SiC nanowires were added to the 6061Al matrix. However, up to now, little work has been reported on the synthesis and properties of insitu forming SiC nanowires in C/CeUHTC composites. Precursor infiltration pyrolysis (PIP) is an effective method to synthesize ceramics in the porous preform due to the ceramic distribution uniformity and the ceramic component designability [27]. In our previous work [21], SiC nanowires were successfully insitu grown on the carbon fiber by PIP method, which distributed inside the C/C performs uniformly. It is worth attempting to use the in-situ grown SiC nanowires to improve the strength and toughness of the C/CeUHTC composites. In this work, the in-situ grown SiC nanowires were introduced into C/C-UHTC composites by PIP method, using the mixture solution of polycarbosilane (PCS), organic HfC polymer (PHC), organic ZrB2 polymer (PZC) as precursor, and ferrocene as catalyst. The effects of in-situ grown SiC nanowires on the microstructure of the ceramic matrix, mechanical property, fracture toughness and particle erosion of the SiC nanowires (NWs) -reinforced C/C -HfC-ZrB2SiC (HZS) were investigated.
2. Experimental procedure 2.1. Materials preparation 2.5-D needle punched carbon fiber fabric (0.45 g/cm3) was densified through thermal gradient chemical vapor infiltration (TCVI) process to produce porous C/C composites with a density of 1.3 g/cm3. Polycarbosilane (PCS), organic HfC precursor (PHC) and ferrocene ((Fe(C5H5)2), purity: 99.0%) were mixed with a weight
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ratio of 1:2:0.01, and then were dispersed in dimethylbenzene to form a homogenous mixed precursor by ultrasonic dispersion at room temperature for 10 h. Fig. 1 shows the schematic preparation procedures of the SiCNWs-C/C composites by PIP method. Firstly, the mix precursor was infiltrated into the porous C/C composites, and then the C/C preforms were dried in a drying oven for 24 h. Finally, they were undergone a heat treatment at 1400 C for 2 h at an argon atmosphere. SiC nanowires were expected to grow in-situ in the C/C preform during the heat treatment. The as received SiCNWsC/C preforms were subjected to densification process by PIP method. The PHC, PCS and organic ZrB2 precursor (POZ) were mixed with a weight ratio of 2:1:1 to form a mixed precursor. After PIP treatment for 10 times, the SiC nanowire-modified C/C-HfCZrB2-SiC composites were fabricated, marked as NW-C/C-HZS composites. Meanwhile, pure C/C preforms without in-situ grown SiC nanowires were subjected to the same PIP treatment under the same conditions for comparisons, marked as C/C-HZS composites.
2.2. Physical and mechanical property tests Density and open porosity of the specimens were measured by Archimedes principle. Each data point was an average of three values. Three-point-bend tests and compression tests were carried out on an electronic universal testing machine (CMT 5304, Suns Co. China) with a loading rate of 0.5 mm/min at room temperature. The sizes of samples used for above tests were 55 mm 10 mm 4 mm. The final results were achieved by computing the average values of at least three specimens. The flexural strength (sf) and flexural modulus (Ef) were calculated from the data recorded during the tests. In addition, the measurements of the fracture toughness (KIC) of as-prepared composites were obtained through three point single edge notched beam (SENB) testing. The initial crack width and length were 200 mm and 2 mm, respectively, which was obtained using blade cutter manually. The particle erosion test was carried out on the shot blasting machine (1010FK, Jichuan Co, China) using Al2O3 particles with a
Fig. 1. Schematic preparation procedures of SiCNWseC/C composites.
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Fig. 2. The XRD patterns of the two modified C/C composites.
diameter of 1~2 mm. The particle flux is 0.24 m3/min. The size of the samples was 430 mm 10 mm. The inner diameter of the gun tip was 5 mm and the distance between the gun tip and the samples was 20 mm. Linear and mass erosion rates were calculated by the change of thickness and weight before and after erosion test. The ultimate result was the average value of three samples. 2.3. Microstructure characterization The morphology and elements of the samples were analyzed using a field emission scanning electron microscopy (FESEM, ZEISSSUPRA 55, Germany), equipped with energy dispersive X-ray spectroscopy (EDS). The phase compositions of the samples were analyzed by X-ray diffraction (XRD, X'Pert PRO MPD) using Cu as anode material. The microstructures of the as-prepared SiC nanowires was characterized by a transmission electron microscope (TEM), high-resolution transmission electron microscopy (HRTEM) and selected area electron diffraction (SAED) images (FEI TecnaiF30G2, USA). Fourier transform infrared (FT-IR) and Raman shift spectrum at room temperature were recorded by a FT-IR spectrometer (Bruker Tensor-27, Germany), a Raman spectrometer (RMS; Renishaw, UK), respectively. 3. Results and discussion 3.1. Microstructure
diffraction peaks. C, SiC, HfC and ZrB2 exist in both samples, as well as HfO2 and SiO2. The existence of HfO2 and SiO2 could be the result of the incomplete pyrolyzation of the precursor under high temperature heat treatment process. Moreover, the diffraction peaks of NW-C/C-HZS is sharper than those of C/C-HZS, indicating that the addition of SiC nanowires could effectively improve the crystallinity of the ceramics. The distribution and morphology of SiC nanowires in C/C performs will directly affect the subsequent ceramic densification process and the properties of the obtained composites. Fig. 3a shows the SEM image of nanowires inside C/C performs. The distribution of nanowires is uniform inside the preform, attributing to the good infiltration ability of the mixed liquid precursor in porous C/C preform. The nanowires are generally several tens to more than one hundred micrometers in length and randomly oriented with straight morphologies, as confirmed by the higher-magnification SEM image in Fig. 3b. To confirm the structure of the nanowires, FT-IR measurement and Raman scattering were employed to analyze the samples. The typical FT-IR spectrum and Raman shift spectrum are shown in Fig. 4. The absorption peaks at around 810, 1080 and 1610 cm 1 were detected by FT-IR. The strong absorption peak at around 810 cm 1 is stretching vibration of Si-C, indicating that the assynthesized nanowires should be SiC. The other two weak absorption bands centered at 1080 and 1610 cm 1 are attributed to the Si-O and H-O-H bending vibration, which might be caused by the incomplete conversion of the precursors and the adsorbed water. Similarly, the Raman shift spectrum shows that the sharp peaks at 790 cm 1 correspond to the modes of transverse optical (TO) phonons of the 3C-SiC [28] in Fig. 4b. In addition, a down shift appears compared to the previous result [29] (TO mode at 796 cm 1 obtained from bulk 3CeSiC). This should be attributed to the confinement effect and inner stress from the structural defect of achieved nanowires. The low-resolution TEM image of the nanowire is shown in Fig. 5a, which shows that the SiC nanowire is linear with uniform structure of about 150 nm in diameter. The corresponding EDS patterns inserted in Fig. 5a shows that the nanowire is composed of Si and C. From selected area electron diffraction (SAED) pattern (inserted in Fig. 5b) and high-resolution TEM (HRTEM) image (Fig. 5b) of a single SiC nanowire, it can be seen that the as-received SiC nanowire has perfect single crystalline structure, possessing high-density stacking faults in the crystal plane normal to the axis of the nanowire. The formation of these stacking faults could consume lower energy compared with the perfect crystal planes [30]. The nanowire crystal lattice fringe spacing is 0.25 nm, which is the same as the distance between {001} planes in 3C-SiC crystalline.
Fig. 2 shows the XRD patterns of the samples. It can be seen that two kinds of modified C/C composites as prepared have similar
Fig. 3. SEM images of nanowires inside C/C performs.
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Fig. 4. FT-IR absorbance of the as-synthesized nanowires (a) and Raman shift spectrum (b).
Fig. 5. TEM image of the single SiC nanowire with the insert showing the corresponding EDS pattern (a) and HRTEM image of the single SiC nanowire and its corresponding SAED pattern (b).
3.2. Mechanical properties Table 1 shows the density, porosity and mechanical properties of the composites. The value of density and porosity are the average value of three specimens measured by Archimedes principle. The C/ C-HZS composites have an average density of 2.27 ± 0.05 g/cm3, with an average porosity of 12.3 ± 0.4%. The NW-C/C-HZS composites have similar density (2.23 ± 0.06 g/cm3) and lower porosity (11.4 ± 0.3%). The obtained flexural strength and fracture toughness are also summarized in Table 1. The average flexural and fracture toughness of the NW-C/C-HZS composites are up to 172 ± 8 MPa and 9.7 ± 0.4 MPa$m1/2, respectively. Compared to the C/C-HZS composites, the flexural and fracture toughness are enhanced by 21% and 28%, respectively. The results show that the incorporation of SiC nanowires greatly improves the flexural strength and fracture toughness of C/C composites modified by HfC-ZrB2-SiC ceramics. Fig. 6 shows the flexural stress-strain curves of the composites. It is obvious that the curves of two composites exhibit a similar trend. It can be seen that in the initial stage of loading load, both materials have a quasi-elastic strain region, and then the flexural strength rapidly declines after the strain reached the maximum
value. It is noteworthy that the curve of the NW-C/C-HZS composites appears a shock region when the strain reaches the maximum value, which is the result of the alternate effects of interfacial debonding and carbon fiber bridging in ceramic matrix, while the C/C-HZS composites break immediately when strain reach the maximum value. In addition, the curve of the C/C-HZS composites shows some small inflection points during the quasiplastic deformation period, which indicates that microcracks appeared in the ceramic matrix, while the curve of the NW-C/CHZS composites is smoother under the same strain. As shown in Fig. 7, the fracture surfaces of the fabricated composites after three point bending test were observed by SEM. In Fig. 7a, the fracture surface is rather planar and few fibers are pulled out, with obvious brittle fracture characteristics, indicating the poor toughness of the C/C-HZS composites. Fig. 7b shows that the carbon fiber bundles are pulled out from the NW-C/C-HZS composites. There are a large number of SiC nanowires and ceramic particles on the surface of fiber bundles, which indicates that the interface between carbon fiber bundles and ceramic matrix was debonding during the process of carbon fiber bundle pull-out and a large amount of ceramic particles remain on the fiber bundle
Table 1 Density, porosity and mechanical properties of composites.
C/C-HZS NW-C/C-HZS
Density (g/cm3)
Porosity (%)
Flexural Strength (MPa)
Fracture Toughness (MPa$m1/2)
Flexural Modulus (GPa)
2.27 ± 0.05 2.23 ± 0.06
12.3 ± 0.4 11.4 ± 0.3
143 ± 9 172 ± 8
7.6 ± 0.3 9.7 ± 0.4
14.2 ± 0.4 18.3 ± 0.3
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Fig. 6. The flexural stress-strain curves of C/C-HZS (a) and NWeC/C-HZS (b) composites.
surface due to the connection of SiC nanowires. Apparently, the fracture surfaces of these two kinds of modified C/C composites are different. The interfacial analysis of the carbon fiber and ceramic matrix was carried out to further understand the cause of fracture and the effect of SiC nanowires on the modified composites. Carbon fiber bundles of the C/C-HZS and NW-C/C-HZS composites in the fracture surface are shown in Fig. 8a and b, respectively. From Fig. 8a, it can be seen clearly that the carbon fiber is closely bonded with the matrix ceramic, and the matrix consists of bulk ceramic particles. A lot of obvious cracks appear in the ceramic matrix, which is due to decomposition and volume
shrinkage of the precursor during the precursor pyrolysis process. From Fig. 8b, the ceramic particles of the NW-C/C-HZS composites are small and uniformly distributed inside the network of SiC nanowires, and the ceramic particles are interconnected by SiC nanowires so that the matrix becomes a unitary body. Compared with C/C-HZS composites, there are a lot of micro holes in the ceramic matrix instead of obvious cracks, due to the existence of SiC nanowires. This kind of structure might cause a weaker interface between the carbon fiber and ceramic matrix. During the fracture process of composite materials, when it is subjected to a certain loading, some cracks will inevitably be formed in the composites, and then the cracks will gradually extend till they cut through the specimen, finally the composite will completely be broken. Therefore, the fracture process of modified C/C composites is actually the process of crack propagation inside the material. There are two general paths for crack propagation: perpendicular to the fiber direction and parallel to the fiber direction. For C/C-HZS composites, due to the strong interfacial bonding between carbon fiber bundles and ceramic matrix, the cracks can not be deflected when the load is transferred to the ceramicpyrolytic carbon interface, and the stress is concentrated at the interface. When the stress gradually increases to the maximum value, the cracks expand along the interface (parallel to the fiber direction). Perpendicular to the fiber direction, the cracks pass directly through the pyrolytic carbon to the fiber, which leads to brittle fracture of pyrolytic carbon matrix and fibers (Fig. 9a). For NW-C/C-HZS composites, the interfacial bonding between the ceramic matrix and the pyrolytic carbon is slightly weakened due to the introduction of SiC nanowires. When the crack propagates to the fiber, the interface starts to debond to prevent the crack from propagating forward. The crack is forced to change its original prolongation direction and extend along the interface by the effect of stress concentration, which eventually causes the debonding of the ceramic matrix and pyrolytic carbon, increasing the fracture
Fig. 7. The fracture surfaces of the fabricated composites after three point bending test. (a) C/C-HZS composites; (b) NW-C/C-HZS composites.
Fig. 8. The interface of the carbon fiber and ceramic matrix of the composites in the fracture surface of C/C-HZS composites (a) and NW-C/C-HZS composites (b).
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Fig. 9. The cracks propagation of the C/C-HZS composites (a) and NW-C/C-HZS composites (b); the cracks propagation in the ceramic matrix of the C/C-HZS composites (c) and the NW-C/C-HZS composites (d).
stress. From Fig. 9b, it is observed that the pyrolytic carbon layer appears delamination phenomenon and the stepped layer is pulled out. Due to the debonding of ceramic matrix and pyrolytic carbon layer, the stress could be released partially. As the cracks encounter obstacles again, the stress is concentrated at the interface between the carbon fiber and the pyrolytic carbon layer since carbon fiber could still bear the load. The mismatch of the coefficient of thermal expansion between the carbon fiber and the pyrolytic carbon layer will result in the generation of circumferential microcracks. The propagation path of circumferential microcracks might be changed along the interface direction when the crack tips are distributed vertically at circumferential microcracks, resulting in the debonding of the interface between the carbon fiber and the pyrolytic carbon layer and the delamination of the pyrolytic carbon layer under the effect of stress. Finally, the carbon fibers break, and they are pulled out from the pyrolytic carbon layer. From Fig. 9c, the ceramic matrix breaks directly because of the low fracture toughness of the ceramics. While the path of the crack propagation is deflected due to the obstacle of the SiC nanowires' network system and ceramic particles' refinement (Fig. 9d). In the ceramic matrix of
the NW-C/C-HZS composites, the ceramic particles are refined because of the introduction of SiC nanowires, and ceramic particles are uniformly distributed in the network system of SiC nanowires and have a more solid connection. When the cracks propagate to the ceramic matrix, the refinement of the ceramic particles can product more obstacles for the crack propagation process, resulting in the deflection of the path of cracks propagation, and the fracture of the SiC nanowires and the deflection of the crack will dissipate extra energy. The effect of SiC nanowires in the fracture process of the NW-C/ C-HZS composites was further investigated. As shown in Fig. 10a, it is clearly observed that there are a large number of SiC nanowires in the interstices of the ceramic particles, which are connected to each other, and ceramic particles are also connected by SiC nanowires. Under the connection of the SiC nanowires, the ceramic matrix becomes an interconnected whole, which enhances the strength of the matrix, thus the mechanical properties of the composites are improved as well. In addition, SiC nanowires formed in situ in the NW-C/C-HZS composites appear in the ceramic blocks after the multiple densification process, which improves the toughness of
Fig. 10. A representative image of ceramic particles bridging by the nanowire (a) numerous pullout or debonding nanowires (b).
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Table 2 Density, porosity and particle erosion performance of composites.
C/C-HZS NW-C/C-HZS
Density (g/cm3)
Porosity (%)
Linear erosion rate (10
2.41 ± 0.05 2.35 ± 0.04
11.6 ± 0.8 12.5 ± 0.6
28 ± 5 11 ± 3
3
mm/s)
Mass erosion rate (10
4
g/s)
3.6 ± 0.6 2.3 ± 0.4
Fig. 11. SEM images of the surface of composites after particle erosion. (a) C/C-HZS composites; (b) NW-C/C-HZS composites; (c) the holes of ceramic matrix of the C/C-HZS composites; (d) the craters of ceramic matrix of the NW-C/C-HZS composites.
ceramic. From Fig. 10b, the pulling out of SiC nanowires can be observed on the fracture surface of ceramic blocks. The stress increases and the microcracks expand when the matrix is under load. Due to the existence of SiC nanowires in some of the ceramic blocks, the propagation of cracks through the ceramic are impeded. The SiC nanowires debond from the ceramic and then pull out, which consumes much energy, so that the ceramic blocks will not crack directly. Therefore, SiC nanowires have a toughening effect on the matrix in this process.
3.3. The particle erosion performance of the composites Table 2 shows the density and the porosity of the two composites after subjected to Al2O3 particle erosion. The C/C-HZS composites have an average density of 2.41 ± 0.05 g/cm3, which corresponds to a porosity of 11.6 ± 0.8%. The composites with SiC nanowires have similar density (2.35 ± 0.04 g/cm3) and similar porosity (12.5 ± 0.6%) with those without SiC nanowires. It shows that the NW-C/C-HZS composites have lower linear and mass erosion rates than the C/C-HZS composites. It is meaningful to investigate the erosion surfaces of different composites to understand the effect of SiC nanowires on the particle erosion performance. Fig. 11 shows the morphology of two modified C/C composites after particle erosion. As shown in Fig. 11a, the surface of the C/C-HZS composites not only has many craters, but also leaves a lot of holes. By further magnification to observe the holes (Fig. 11c), it is clear to see that the Al2O3 particles directly puncture the short-cut fiber web of the ceramic matrix, and the next non-woven layer can be seen through the hole. Comparatively,
the surface of the NW-C/C-HZS composites also leaves many craters, but just few pore after particle erosion (Fig. 11b). Fig. 11d shows the magnification SEM image of the crater of the NW-C/CHZS composite. Compared to the craters of the C/C-HZS composites, the size of the craters of the NW-C/C-HZS composites is smaller, and the short-cut fiber web is not punctured by high energy Al2O3 particles. To further investigate the erosion morphologies of two modified composites, the craters were observed by SEM. Fig. 12 shows the craters on the ceramic matrix of two modified C/C composites after particle erosion. As shown in Fig. 12a, the big bulk ceramic particles in the ceramic matrix of the C/C-HZS composite are crushed and peeled off. From Fig. 12b, it is clear that the size of ceramic particles of the NW-C/C-HZS composite is much smaller than that of the C/CHZS composite. From Fig. 12c and d, the size of ceramic particles of the NW-C/C-HZS composites is about 1e4 mm and three kinds of ceramic particles are uniformly distributed, embedded in each other. In addition, SiC nanowires can also be found between the ceramic particles in Fig. 12d, which connect the ceramic particles into a whole. During the particle erosion process, ceramic particles can effectively transfer energy to each other through SiC nanowires, so that the ceramic matrix has better erosion resistance.
4. Conclusions The SiC nanowires reinforced C/C-HfC-ZrB2-SiC composites were fabricated by a multistep technique of precursor infiltration and pyrolysis process. Due to the introduction of SiC nanowires, the size of the ultra-high temperature ceramic particles were refined
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Fig. 12. The craters leave on the ceramic matrix after particle erosion of C/C-HZS composites (a) and NW-C/C-HZS composites (b); (c) and (d): high resolution images of (b).
greatly, and the ceramic matrix becomes a unitary body with the interconnection of SiC nanowires, hence the NW-C/C-HZS composites displayed good particle erosion resistance. The flexural strength, flexural modulus and fracture toughness of the NW-C/CHZS composites increased by 21%, 29% and 28%, respectively compared to the C/C-HZS composites without SiC nanowires. The improvement of toughness and strength of the HZS ceramics is attributed to the pullout of carbon fiber, the deflection of cracks induced by nanowires and the pullout of SiC nanowires. Acknowledgments This work was supported by National Natural Science Foundation of China [grant numbers 51521061, 51432008 and 51572223] and the “111” Project [grant number B08040]. References [1] X. Shen, K. Li, H. Li, et al., Microstructure and ablation properties of zirconium carbide doped carbon/carbon composites, Carbon 48 (2) (2010) 344e351. [2] L. Gaab, D. Koch, G. Grathwohl, Effects of thermal and thermomechanical induced mechanical changes of C/C composites, Carbon 48 (10) (2010) 2980e2988. [3] J.C. Rietsch, J. Dentzer, A. Dufour, F. Schnell, L. Vidal, P. Jacque, et al., Characterizations of C/C composites and wear debris after heavy braking demands, Carbon 47 (1) (2009) 85e93. [4] Y.L. Zhang, Z. Hu, B. Yang, et al., Effect of pre-oxidation on the ablation resistance of ZrB 2 eSiC coating for SiC-coated carbon/carbon composites, Ceram. Int. 41 (2) (2015) 2582e2589. [5 ] T.L. Dhami, O.P. Bahl, B.R. Awasthy, Oxidation-resistant carbon-carbon composites up to 1700 C, Carbon 33 (4) (1995) 479e490. [6] Q.G. Fu, L. Zhuang, Q.W. Ren, et al., Carbon nanotube-toughened interlocking buffer layer to improve the adhesion strength and thermal shock resistance of SiC coating for C/CeZrCeSiC composites, J. Materiomics 1 (3) (2015) 245e252. [7] D.D. Jayaseelan, R.G.D. S a, P. Brown, et al., Reactive infiltration processing (RIP) of ultra high temperature ceramics (UHTC) into porous C/C composite tubes, J. Eur. Ceram. Soc. 31 (3) (2011) 361e368. [8] L. Liu, H. Li, W. Feng, et al., Effect of surface ablation products on the ablation resistance of C/CeSiC composites under oxyacetylene torch, Corros. Sci. 67 (2013) 60e66. [9] K.Z. Li, X. Jing, Q.G. Fu, et al., Effects of porous C/C density on the densification behavior and ablation property of C/CeZrCeSiC composites, Carbon 57 (3) (2013) 161e168.
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