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carbon composites by a novel SiC nanoribbon-reinforced SiC–Si ceramic coating
Corrosion Science 92 (2015) 272–279
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Corrosion Science journal homepage: www.elsevier.com/locate/corsci
Oxidation protection of carbon/carbon composites by a novel SiC nanoribbon-reinforced SiC–Si ceramic coating Yanhui Chu a,b, Hejun Li a,⇑, Huijuan Luo a, Lu Li a, Lehua Qi c a
State Key Laboratory of Solidification Processing, Carbon/Carbon Composites Research Center, Northwestern Polytechnical University, Xi’an 710072, China School of Engineering and Applied Sciences, Harvard University, Cambridge, MA 02138, USA c School of Mechatronic Engineering, Northwestern Polytechnical University, Xi’an 710072, China b
a r t i c l e
i n f o
Article history: Received 5 July 2014 Accepted 17 December 2014 Available online 24 December 2014 Keywords: A. Ceramic matrix composites C. High temperature corrosion C. Oxidation
a b s t r a c t A novel SiC–Si ceramic coating reinforced with well-dispersed SiC nanoribbons was prepared on carbon/ carbon (C/C) composites by chemical vapor deposition and pack cementation. The mechanical tests showed that the fracture toughness of the coating and the interfacial bonding strength of the coatingC/C increased by 98% and 99% after incorporating SiC nanoribbons, respectively. The oxidation tests showed that the oxidation of the samples was a continuous weight gain process, and the final weight gain were 61.84 gm2 during thermogravimetric analysis from 300 °C to 1500 °C and 38.0 gm2 after isothermal oxidation test for 96 h, respectively. Ó 2014 Elsevier Ltd. All rights reserved.
1. Introduction Oxidation protection is an essential requirement for carbon/carbon (C/C) composites to be used as high-temperature structural materials in an oxygen-containing environment [1]. An efficient anti-oxidation coating is a key technique for addressing this issue [2]. Silicon-based ceramic coatings have been attracted extensive attention due to their good oxidation resistance and compatibility with C/C composites [3–5]. The common failure mode of these ceramic coatings is cracking during cool-down to low temperatures, which limits their capability for protecting of C/C composites in a large temperature range. Toward solving this problem, SiC nanowires have been incorporated into these coatings as the reinforcing materials [6–8], which played a positive role. However, the toughening effects of SiC nanowires in the coatings are not very exciting [9]. Compared to the nanowires, the nanoribbons with a special nanostructure are considered as more effective reinforcing materials in the composites because they can establish the good connectivity with the matrix. The two main factors responsible for this are as follows [10,11]: (1) Surface area: The nanofiller with a sheet or belt-like geometry has larger interfacial contact area with the matrix materials than that of the nanofiller with a cylinder geometry and the same cross-sectional area. (2) Geometry: Compared with the nanofiller with cylinder geometry, it is easier for the matrix materials to adhere to the nanofiller with a sheet
⇑ Corresponding author. Tel.: +86 29 88495004; fax: +86 29 88492642. E-mail address: [email protected] (H. Li). http://dx.doi.org/10.1016/j.corsci.2014.12.013 0010-938X/Ó 2014 Elsevier Ltd. All rights reserved.
or belt-like geometry. In our previous work, SiC nanoribbons have been synthesized in large-scale using ferrocene as catalyst by a new combination growth mechanism of vapor–liquid–solid-based and vapor–solid-tip [12,13], unlike previously reported one-step vapor–solid growth for SiC nanowires [6] or two-step vapor–solid growth for bamboo-shaped SiC nanowires [14,15]. Provided that SiC nanoribbons are designed to incorporate into the silicon-based ceramic coatings, such as SiC–Si, it can efficiently suppress the cracking of the coatings, which may improve the oxidation inhibition ability of the coated C/C composites. In the present work, a novel SiC–Si ceramic coating reinforced with well-dispersed SiC nanoribbons was designed and prepared on C/C composites by a two-step technique involving chemical vapor deposition (CVD) and pack cementation. Firstly, a porous SiC nanoribbon layer was in-situ grown on C/C composites by CVD. Secondly, this material was subjected to pack cementation to infiltrate the porous layer with the coating materials. The microstructure, mechanical and oxidation resistant properties of the coated C/C samples were investigated, as well as the reinforcement mechanism of SiC nanoribbons. 2. Experimental procedure Cubic samples (8 mm 8 mm 8 mm) as substrates, which were cut from bulk two-dimensional C/C composites with a density of 1.72 g/cm3, were hand-abraded using SiC grit paper (300 mesh), then cleaned ultrasonically with ethanol and dried at 100 °C for 2 h. SiC nanoribbon-reinforced SiC–Si ceramic coating was prepared by the following two-step technique of CVD and pack
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cementation. The first step technique of CVD for synthesizing the porous SiC nanoribbon layer on C/C samples is described as follows: The chemical composition of precursor powders was given as follows: 40–70% SiO2 (300 mesh), 10–15% Si (300 mesh), 15–30% graphite (320 mesh), and 5–15% ferrocene (300 mesh). The above powders were weighed up, mixed by tumbling in a ball mill for 2–4 h, and then put into a graphite crucible. C/C samples were suspended over the mixture powder in the graphite crucible with graphite lid and placed into the center of a vertical graphite furnace. The furnace was heated to 1700–1800 °C at the rate of 10 °C/min in argon and then maintained at that temperature for 2–4 h. After heating, the furnace was naturally cooled down to room temperature. The precursor powder of SiC–Si ceramic coating for the second step technique of pack cementation was mixed as follows: 70–85 wt.% Si (300 mesh), 5–15 wt.% SiC (300 mesh), 7–15 wt.% graphite (320 mesh) and 2–15 wt.% Al2O3 (300 mesh). SiC nanoribbon-C/C samples were immersed in above mixtures in the graphite crucible. Then the whole assembly was heat-treated in argon at 1900–2100 °C for 2 h to form the desire SiC nanoribbon-reinforced SiC–Si ceramic coating on C/C samples. Mechanical properties of the coatings were measured by indentation method. Micro-indentation test was performed using Nanovea universal-indenter (Nanovea Inc., USA) with a diamond Berkovich indenter to measure the hardness and elastic modulus of the coatings. Before measurement, microindentation set up was calibrated using a standard steel sample. After that the indentations were carried out on the well-polished cross-section of the coatings at load of 0.5 N. The load was applied linearly up to the maximum load in 60 s followed by unloading in 60 s. The hardness and elastic modulus were also calculated by the Oliver and Pharr method [16]. A total of 10 indentations with the same load were carried out on each sample. Fracture toughness (Kc) of the coatings was measured on the well-polished surface of the coatings by the indentation technique using the same micro-indentation tester with a Vickers indenter. The loads with 5 N and 49 N were used to generate the cracks in the coatings for the calculation of the fracture toughness using the following equation [17]
2=5 E P Kc ¼ v 1=2 H al
273
consisting of oxygen and argon with a volume ratio of 22:78 because the initial oxidation temperature of C/C composites was about 370 °C [18]. Weight changes of the coated C/C samples related to their temperature were recorded with thermogravimetric mode. Isothermal oxidation test of the coated C/C samples was carried out at 1500 °C in the corundum tube furnace with natural convection of air to investigate their isothermal oxidation behavior. After the furnace was heated up to 1500 °C, the coated C/C samples were put directly into the furnace, whereafter they were taken out at the designated time and cooled at room temperature in air for weighing. After that they were put into the furnace again for the next isothermal oxidation test. An electronic balance with a sensitivity of ±0.1 mg was used to calculate the weight change (W) of the samples by the following equation:
W¼
m1 m0 S
ð3Þ
where m0 and m1 are the weight of the coated samples before and after oxidation, respectively, S is the surface area of the coated samples. The morphology and microstructure of the samples were analyzed by a scanning electron microscopy (SEM, Quanta 600FEG, FEI, USA), equipped with energy dispersive X-ray spectroscopy (EDS). 3. Results and discussion Fig. 1(a) is a typical SEM surface image showing the morphology of the porous SiC nanoribbon layer on C/C composites. It is obvious that the porous layer exhibits an open and porous structure consisting of a large quantity of wire-like and curvature nanostructures with typical lengths in the range of several tens to several hundreds of micrometers; some of them even have lengths on the order of millimeters. From high-magnification SEM images
ð1Þ
where P is the applied load, E is the elastic modulus, H is the hardness, a is the half diagonal of the indents, l is the radial crack length, and v is the constant. In this study, v = 0.0122 because l/a ratios were between 1 and 2.5. A total of 5 indentations were performed on each sample. Tensile test was carried out using an Instron universal test machine (CMT5304-30 kN, Instrument Co., Ltd., China) to estimate the interfacial bonding strength between the coating and C/C substrate (ASTM standard C633). The cylin-drical aluminum rods with dimensions of U25 mm 25 mm were used as matching parts. The coated C/C samples were bonded with the matching parts at both ends by epoxy resins. After being solidified for 72 h at room temperature, they were tested by the universal test machine. The maximum force of each sample was recorded. The bonding strength (r) is calculated using the following equation:
r¼
F S
ð2Þ
where F is the maximum force and S is the cross-sectional area of the samples. A total of 3 samples were measured for each group. Thermogravimetric analysis (TGA) of the coated C/C samples was performed using a Metter Toledo Star TGA/SDTA 851 thermal analyzer to investigate their oxidation behavior between room temperature to 1500 °C. The coated C/C samples were heated from 300 °C to 1500 °C with a heating rate of 20 °C/min in simulated air
Fig. 1. (a) SEM surface image of the porous SiC nanoribbon layer on C/C composites; (b) high-magnification image of (a).
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Fig. 2. (a) SEM surface image of the coating with SiC nanoribbons; (b) cross-section backscattering electron image of the coating with SiC nanoribbons.
(Fig. 1(b)), a high width-to-thickness ratio of the nanoribbons can be clearly observed. The nanoribbons have smooth surfaces along their length and width and most of them have uniform rectangular cross sections with depths of 10–100 nm and widths of 100– 900 nm. Fig. 2(a) shows that SEM surface image of the coating with SiC nanoribbons. Obviously, the coating exhibited a dense surface morphology, on which the free Si was well distributed around SiC particles and a small quantity of Al2O3 was observed. Cross-section backscattering electron image (Fig. 2(b)) displays that the coating had a dense microstructure with the major SiC phase and Si phase, in which Si phase was well filled into the gaps among SiC phase and few holes and penetrating cracks were observed. SEM images shown in Fig. 3 display the morphology of the fracture surface of SiC nanoribbon-reinforced SiC–Si coating. From Fig. 3(a), the morphology of the fracture surface of the coating shows clearly a stepwise feature, indicating the crack occurred an obvious deflection during the coating failure. Of particular interest is that there are numerous SiC nanoribbons on the fracture surface. Most of SiC nanoribbons are observed to debond from the matrix, as exemplified in Fig. 3(b). These nanoribbons are partially exposed to the crack plane, and the exposed area of such nanoribbons lies parallel to the crack plane. Therefore, such nanoribbons induce the deflection of the crack during crack propagation, leading to the formation of the stepwise fracture surface. In addition, note that some SiC nanoribbons exhibit an obvious pullout feature on the fracture surface, which can absorb vast amounts of fracture energy due to the presence of the good connectivity between the nanoribbons and coating matrix. Fig. 3(c) is a typical SEM image showing bridging feature of SiC nanoribbon within the crack that resists the crack widening (red1 arrow in Fig. 3(c)), suggesting a 1 For interpretation of color in Figs. 3–5, 7 and 8, the reader is referred to the web version of this article.
potential crack bridge toughening mechanism. Moreover, Fig. 3(c) also shows that the crack propagates to SiC nanoribbon and then deflects to another direction (blue arrow in Fig. 3(c)), confirming that there is indeed the toughening mechanism of crack deflection in this coating. As shown in Fig. 3(d), the cracks showed markedly deflection near the nanoribbons when propagating near the nanoribbons, and then propagated along the nanoribbon boundary. Due to the good connectivity between the matrix and the nanoribbons, the propagation of crack would be resisted and turned to another direction when propagating along the nanoribbon boundary to a certain degree (red arrows in Fig. 3(d)). This process results in an increase in the propagation path and resistance of the crack, which can cause the coating to be toughened. The elastic modulus and hardness values of the coatings obtained from the micro-indentation tests are shown in Table 1. It can be found that the hardness and elastic modulus of the coating increase by 27% and 39% with the incorporation of SiC nanoribbons, respectively. The fracture toughness of the coatings was also measured by micro-indentation technology. First, the fracture toughness measurement at a low load was evaluated on the cross-section of the coatings to avoid substrate effect. Fig. 4(a) shows SEM image of the indentation generated at 5 N load on the cross-section of the coating without SiC nanoribbons. The radial cracks which propagated from the indentation corner, where the highest stresses occurred, were observed. According to the above results, the fracture toughness of the coatings can be calculated using Eq. (1). The result is presented in Table 1. The coating without SiC nanoribbons shows an average fracture toughness of 7.28 ± 0.85 MPam1/2. SEM image of the indentation generated at 5 N load on the cross-section of the coating with SiC nanoribbons was shown in the left-bottom of Fig. 4(b). An indentation with the diagonal length of 20 lm was observed, but no radial crack was found around the indentation. Considering the thickness of the coating has only 85 lm, the fracture toughness measurement at a high load was evaluated on the surface of the coatings. SEM image of the indentation generated at 49 N load on the surface of the coating with SiC nanoribbons is presented in Fig. 4(b). It was found that the diagonal length of the indentation was about 100 lm, which was larger than the thickness of the coating. It should be noted that the radial cracks propagated from the indentation corner were observed, and the fracture toughness of the coatings was calculated using the Eq. (1), which was presented in Table 1. The coating with SiC nanoribbons shows an average fracture toughness of 14.42 ± 0.97 MPam1/2. An increase in 98% in relative fracture toughness of the coating was observed with the incorporation of SiC nanoribbons. The reinforcement effectiveness of SiC nanoribbons is better than the reported results about SiC nanowires in our previous work [6,8,9]. Detailed SEM investigations inside the indentation cracks reveal the following factors responsible for the improvement in fracture toughness of the coatings. From Fig. 4(b), it is interesting to find that the radial cracks exhibit the zigzag propagation path in the coating. High-magnification SEM image (Fig. 4(c)) shows that the crack exhibits obvious deflection feature when it encounters SiC nanoribbons (red arrow in Fig. 4(c)) during its propagation, which results in the formation of the zigzag propagation path and finally contributes to an increase in the toughness of the coating. Furthermore, Fig. 4(d) shows the high-magnification image within the crack which clearly shows numerous nanoribbon pullout and bridged features within the crack that resist the crack widening (red arrow in Fig. 4(d)) and hence contribute toward the toughening of the coating. In addition to the fracture toughness of the coatings, the interfacial bonding between the coatings and C/C composites is also regarded as the key factor to determine the oxidation resistant capacity of the coated C/C composites. The good interfacial bonding can alleviate thermal stress between the coatings and C/C composites, which ultimately affects the oxidation resistant
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Fig. 3. SEM fracture surface images of SiC nanoribbon-reinforced SiC–Si coating: (a) a representative SEM image of numerous debonding and pullout nanoribbons; (b) a representative SEM image of numerous debonding nanoribbons (magnified image of (a)); (c) a representative SEM image of crack bridging by the nanoribbon and crack deflection near the nanoribbon; (d) a representative SEM image of crack deflection near the nanoribbon.
Table 1 Micro-mechanical properties of the coatings. Samples
Hardness (GPa)
Elastic modulus (GPa)
Fracture toughness (MPam1/2)
Bonding strength (MPa)
Coating Nanoribbon-coating
24.53 ± 0.64 31.04 ± 0.46
258.89 ± 7.39 358.58 ± 11.37
7.28 ± 0.85 14.42 ± 0.97
21.53 ± 1.61 42.86 ± 2.54
Fig. 4. (a) SEM image of 5 N Vickers indentations on the cross-section of the coating without SiC nanoribbons; (b) SEM image of 49 N Vickers indentations on the surface of the coating with SiC nanoribbons (the inset is SEM image of 5 N Vickers indentations on the cross-section of the coating with SiC nanoribbons); (c) the representative crack deflection feature near the nanoribbons (enlarged image of the area marked as ‘‘A’’ in (b)); (d) the representative nanoribbon pullout and bridged features within the crack (enlarged image of the area marked as ‘‘B’’ in (b)).
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capacity of the coated C/C composites [19]. Table 1 presents a comparative study of the interfacial bonding strength between C/C composites and coatings. After incorporating SiC nanoribbons, the interfacial bonding strength between the coating and C/C composites represents a 99% increase (from 21.53 ± 1.61 to 42.86 ± 2.54 MPa). SEM fracture surface image of the interfacial area between the coating and C/C substrate is displayed in Fig. 5 (a). Numerous pullout and debonding SiC nanoribbons were observed on the fracture surface. More importantly, some SiC nanoribbon bridging features were detected between the coating and C/C substrate (red arrows in Fig. 5(a)), which implied the presence of interface anchoring effect of SiC nanoribbon between the coating and C/C substrate. A schematic diagram of in situ grown SiC nanoribbons on the interfacial area between the coating and C/C substrate is shown in Fig. 5(b). It is evinced here that some in situ grown SiC nanoribbons in the pores of C/C substrate established the anchoring effect between SiC nanoribbons and C/C substrate, as shown in Fig. 5(c) (high-magnification of Fig. 5(b)). These SiC nanoribbons may attach themselves strongly to the surrounding fibers. During SiC nanoribbon pullout, the relative sliding can occur between SiC nanoribbons and the surrounding fibers, which will resist the pullout of SiC nanoribbons, resulting in an improvement in the interfacial bonding between the coating and C/C substrate. Fig. 6(a) shows that the weight change curve of the coated C/C samples with SiC nanoribbons during TGA test from 300 °C to
1500 °C. It can be seen that the oxidation of the samples was a continuous weight gain process and the final weight gain was 61.84 gm2 during TGA test, which suggested that the as-prepared coating had good oxidation protective ability for C/C composites between room temperature and 1500 °C. The season for this is because the incorporated SiC nanoribbons effectively suppressed the cracking of the coating by via various toughening mechanisms including nanoribbon pullout, microcrack bridging by nanoribbon and microcrack deflection, avoiding the formation of the penetrating cracks in the coating, as shown in Fig. 2(b). The weight gain of the samples was believed to due to the oxidation of the coating, which could be divided into two stages, marked as A and B, respectively. During the stage A (300–1400 °C), the weight gain of the samples was a slow process in that the oxidation rate of SiC and Si phases in the coating was slow as the temperature was lower than 1400 °C. As the temperature was over 1400 °C, the oxidation rate of SiC and Si phases in the coating would be accelerated. Therefore, the weight gain of the samples exhibited a rapid process during the stage B (1400–1500 °C). Fig. 6(b) displays the weight change of the coated C/C samples with SiC nanoribbons during isothermal oxidation test at 1500 °C in air. Clearly, the oxidation of the samples was also a continuous weight gain process during the isothermal oxidation test, which is quite consistent with a parabolic law. As shown in Fig. 6(c), square of the weight gain as a function of oxidation time, the linear relationship confirms that the oxidation follows parabolic behavior, suggesting that the
Fig. 5. (a) SEM fracture surface image of the interfacial area between the coating and C/C substrate; (b) and (c) schematic description of the anchoring mechanism of in situ grown SiC nanoribbons on the interfacial area between the coating and C/C substrate.
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277
Weight gain (g•m-2)
70 60
(a)
50 40 30 20
B
A
10 0 300
600
900
1200
1500
Weight gain (g•m-2)
Temperature/
40
(b)
30 20 10 0 0
20
40
60
80
100
Weight gain (g 2 •m -4 )
Oxidation time (h) 1600 1400 1200 1000 800 600 400 200 0
(c)
Fig. 7. (a) A representative SEM image of nanoribbon pullout within the crack on the surface of the coating with SiC nanoribbons after isothermal oxidation test; (b) SEM cross-section image of coated C/C samples with SiC nanoribbons after isothermal oxidation test (the inset is EDS analysis of spot A in (b)).
0
20
40
60
80
100
Oxidation time (h) Fig. 6. (a) Weight change curve of the coated C/C samples with SiC nanoribbons during TGA test with the increasing temperature from 300 °C to 1500 °C at the rate of 20 °C/min in air (the total oxidation time: 1 h); (b) weight change of the coated C/C samples with SiC nanoribbons as a function of oxidation time during isothermal oxidation test at 1500 °C in air; (c) square of the weight gain of the coated C/C samples with SiC nanoribbons as a function of oxidation time during isothermal oxidation test at 1500 °C in air.
oxidation is controlled by a diffusion process. The parabolic oxidation rate constants, kp, was calculated to be 13.52 gm2h1, less than that of the reported pure SiC ceramic [20], from the slope of the lines according to the following equation
W 2 ¼ kp t
ð4Þ
where W is the weight change of the samples, t the oxidation time. Meanwhile, the weight gain of the samples was 38.0 gm2 after isothermal oxidation at 1500 °C in air for 96 h, which is also less than that that of the samples after TGA test. These suggested that the weight gain of the samples was a combination of the weight gain resulting from the oxidation of the coating and the weight loss resulting from the oxidation of C/C substrate during isothermal oxidation. Obviously, the oxidation of C/C substrate was not dominant but slight and the oxidation of the coating was dominant during isothermal oxidation. Therefore, it can be inferred that the as-obtained coating has the potential to protect C/C composites against oxidation in a large temperature range. Fig. 7(a) shows SEM surface image of the coated C/C samples with SiC nanoribbons after isothermal oxidation test. A smooth SiO2 glass layer was formed on the coating surface (blue arrows
in Fig. 7(a)). Meanwhile, some cracks were observed, in which there are some pullout SiC nanoribbons (red arrow in Fig. 7(a)) and SiO2 glass (blue arrows in Fig. 7(a)), implying that they played a key role in suppressing the cracking of the crack and sealing the crack itself. Fig. 7(b) shows SEM cross-section image of the coated C/C samples with SiC nanoribbons after isothermal oxidation test, on which the following features can be found. First, a SiO2 glass layer was observed on the coating surface, in which there are some perfect bubbles and broken bubbles (red arrow in Fig. 7(b)) due to the oxidation of the coating and C/C substrate [9]. Second, the cracks were detected, which had penetrated the coating and extended into C/C substrate, resulting in a slight oxidized zone without void in C/C substrate. Of particular interest is that the penetrating cracks were not open but sealed with some materials. EDS analysis (Fig. 7(b)) shows that these materials involve Si, O, Al and Pt elements. The presence of Si, O and Al elements confirms that these materials are the generated SiO2 glass [15], indicating that the penetrating cracks can be sealed by the generated SiO2 glass at elevated temperatures. The presence of Pt element can be regarded as an impurity from the spray-gold for SEM analysis. Therefore, it can be inferred that the oxidation of C/C was controlled by the diffusion rate of O2 in the generated glass layer within the penetrating cracks in the coating at elevated temperatures. To explain the oxidation mechanisms of the coated C/C samples with SiC nanoribbons during isothermal oxidation, a schematic diagram is presented in Fig. 8. First, the samples would undergo thermal cycles for 6 times between 1500 °C and room temperature during the isothermal oxidation test for weighing. During this process, the large thermal stress would induce the cracking of the coating due to the mismatch of the coefficient of thermal expansion between the coating and C/C, however, the cracking of the coating
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gases would be accumulated together with the gases generated by the oxidation of the coating to form the bubbles in the generated glass layer, as shown in Fig. 8(b). The bubbles would break as the pressure of these gases exceeded the surface tension of the bubbles. This was also the season why the broken bubbles were found preferentially on the coating surface along the penetrating cracks (red arrow in Fig. 7(b)). It should be noted that the oxidation of C/C substrate is slight during isothermal oxidation in that it is controlled by the diffusion rate of O2 in the generated glass layer. 4. Conclusions In this study, a novel SiC nanoribbon-reinforced SiC–Si ceramic coating had been prepared on C/C composites by a two-step technique consisting of CVD and pack cementation. The as-prepared coating had a dense microstructure, in which SiC nanoribbons were well dispersed. The incorporated SiC nanoribbons significantly improved the hardness, elastic modulus and fracture toughness of the coating, which was primarily attributable to the nanoribbon pullout, nanoribbon bridging, crack deflection. Meantime, the incorporated SiC nanoribbons also greatly improved the interfacial bonding of the coating-substrate, which was mainly resulted from the anchoring effect of SiC nanoribbons at the coating-substrate interface. The as-received coating possessed excellent oxidation protective capacity for C/C composites in a large temperature range. This study provides a promising way to improve the oxidation protective capacity of ceramic coatings for carbon materials by the reinforcement effects of SiC nanoribbons. Acknowledgements This work has been supported by the National Natural Science Foundation of China (Grant Nos. 51221001, 51272212, and 51275417), the ‘‘111’’ Project (Grant No. B08040), the State Scholarship Fund of China Scholarship Council, the Doctorate Foundation of Northwestern Polytechnical University and Excellent Doctorate Foundation of Northwestern Polytechnical University. References
Fig. 8. Schematic description of possible oxidation mechanisms of the coated C/C samples with SiC nanoribbons during isothermal oxidation.
could be effectively suppressed by via various toughening mechanisms of SiC nanoribbons including nanoribbon pullout, microcrack bridging by nanoribbon and microcrack deflection. Thus, the size and number of the cracks in the coating were decreased to a large extent. At the same time, the generated SiO2 glass possessed the good fluidity at elevated temperatures, which could effectively seal the generated penetrating cracks in the coating, as shown in Fig. 8(a). Thus, the as-obtained coating has a good oxidation protective ability for C/C composites. While the presence of Al3+ could damage the network structure of the SiO2 and result in the increase of oxygen diffusion in the generated SiO2 glass [9]. For one thing some oxygen diffused easily to the surface of the coating through the glass layer (indicated by red arrows), continued to react with the coating and generated some gases such as CO, CO2 and SiO (indicated by blue arrows in Fig. 8(b)) [9]. For another, some oxygen also diffused easily to C/C substrate through the generated SiO2 glass within the penetrating cracks in the coating (indicated by red arrows), leading to the oxidation of C/C substrate and releasing CO2 and CO gases (indicated by blue arrows in Fig. 8(b)) [9]. These
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Contents lists available at ScienceDirect
Corrosion Science journal homepage: www.elsevier.com/locate/corsci
Oxidation protection of carbon/carbon composites by a novel SiC nanoribbon-reinforced SiC–Si ceramic coating Yanhui Chu a,b, Hejun Li a,⇑, Huijuan Luo a, Lu Li a, Lehua Qi c a
State Key Laboratory of Solidification Processing, Carbon/Carbon Composites Research Center, Northwestern Polytechnical University, Xi’an 710072, China School of Engineering and Applied Sciences, Harvard University, Cambridge, MA 02138, USA c School of Mechatronic Engineering, Northwestern Polytechnical University, Xi’an 710072, China b
a r t i c l e
i n f o
Article history: Received 5 July 2014 Accepted 17 December 2014 Available online 24 December 2014 Keywords: A. Ceramic matrix composites C. High temperature corrosion C. Oxidation
a b s t r a c t A novel SiC–Si ceramic coating reinforced with well-dispersed SiC nanoribbons was prepared on carbon/ carbon (C/C) composites by chemical vapor deposition and pack cementation. The mechanical tests showed that the fracture toughness of the coating and the interfacial bonding strength of the coatingC/C increased by 98% and 99% after incorporating SiC nanoribbons, respectively. The oxidation tests showed that the oxidation of the samples was a continuous weight gain process, and the final weight gain were 61.84 gm2 during thermogravimetric analysis from 300 °C to 1500 °C and 38.0 gm2 after isothermal oxidation test for 96 h, respectively. Ó 2014 Elsevier Ltd. All rights reserved.
1. Introduction Oxidation protection is an essential requirement for carbon/carbon (C/C) composites to be used as high-temperature structural materials in an oxygen-containing environment [1]. An efficient anti-oxidation coating is a key technique for addressing this issue [2]. Silicon-based ceramic coatings have been attracted extensive attention due to their good oxidation resistance and compatibility with C/C composites [3–5]. The common failure mode of these ceramic coatings is cracking during cool-down to low temperatures, which limits their capability for protecting of C/C composites in a large temperature range. Toward solving this problem, SiC nanowires have been incorporated into these coatings as the reinforcing materials [6–8], which played a positive role. However, the toughening effects of SiC nanowires in the coatings are not very exciting [9]. Compared to the nanowires, the nanoribbons with a special nanostructure are considered as more effective reinforcing materials in the composites because they can establish the good connectivity with the matrix. The two main factors responsible for this are as follows [10,11]: (1) Surface area: The nanofiller with a sheet or belt-like geometry has larger interfacial contact area with the matrix materials than that of the nanofiller with a cylinder geometry and the same cross-sectional area. (2) Geometry: Compared with the nanofiller with cylinder geometry, it is easier for the matrix materials to adhere to the nanofiller with a sheet
⇑ Corresponding author. Tel.: +86 29 88495004; fax: +86 29 88492642. E-mail address: [email protected] (H. Li). http://dx.doi.org/10.1016/j.corsci.2014.12.013 0010-938X/Ó 2014 Elsevier Ltd. All rights reserved.
or belt-like geometry. In our previous work, SiC nanoribbons have been synthesized in large-scale using ferrocene as catalyst by a new combination growth mechanism of vapor–liquid–solid-based and vapor–solid-tip [12,13], unlike previously reported one-step vapor–solid growth for SiC nanowires [6] or two-step vapor–solid growth for bamboo-shaped SiC nanowires [14,15]. Provided that SiC nanoribbons are designed to incorporate into the silicon-based ceramic coatings, such as SiC–Si, it can efficiently suppress the cracking of the coatings, which may improve the oxidation inhibition ability of the coated C/C composites. In the present work, a novel SiC–Si ceramic coating reinforced with well-dispersed SiC nanoribbons was designed and prepared on C/C composites by a two-step technique involving chemical vapor deposition (CVD) and pack cementation. Firstly, a porous SiC nanoribbon layer was in-situ grown on C/C composites by CVD. Secondly, this material was subjected to pack cementation to infiltrate the porous layer with the coating materials. The microstructure, mechanical and oxidation resistant properties of the coated C/C samples were investigated, as well as the reinforcement mechanism of SiC nanoribbons. 2. Experimental procedure Cubic samples (8 mm 8 mm 8 mm) as substrates, which were cut from bulk two-dimensional C/C composites with a density of 1.72 g/cm3, were hand-abraded using SiC grit paper (300 mesh), then cleaned ultrasonically with ethanol and dried at 100 °C for 2 h. SiC nanoribbon-reinforced SiC–Si ceramic coating was prepared by the following two-step technique of CVD and pack
Y. Chu et al. / Corrosion Science 92 (2015) 272–279
cementation. The first step technique of CVD for synthesizing the porous SiC nanoribbon layer on C/C samples is described as follows: The chemical composition of precursor powders was given as follows: 40–70% SiO2 (300 mesh), 10–15% Si (300 mesh), 15–30% graphite (320 mesh), and 5–15% ferrocene (300 mesh). The above powders were weighed up, mixed by tumbling in a ball mill for 2–4 h, and then put into a graphite crucible. C/C samples were suspended over the mixture powder in the graphite crucible with graphite lid and placed into the center of a vertical graphite furnace. The furnace was heated to 1700–1800 °C at the rate of 10 °C/min in argon and then maintained at that temperature for 2–4 h. After heating, the furnace was naturally cooled down to room temperature. The precursor powder of SiC–Si ceramic coating for the second step technique of pack cementation was mixed as follows: 70–85 wt.% Si (300 mesh), 5–15 wt.% SiC (300 mesh), 7–15 wt.% graphite (320 mesh) and 2–15 wt.% Al2O3 (300 mesh). SiC nanoribbon-C/C samples were immersed in above mixtures in the graphite crucible. Then the whole assembly was heat-treated in argon at 1900–2100 °C for 2 h to form the desire SiC nanoribbon-reinforced SiC–Si ceramic coating on C/C samples. Mechanical properties of the coatings were measured by indentation method. Micro-indentation test was performed using Nanovea universal-indenter (Nanovea Inc., USA) with a diamond Berkovich indenter to measure the hardness and elastic modulus of the coatings. Before measurement, microindentation set up was calibrated using a standard steel sample. After that the indentations were carried out on the well-polished cross-section of the coatings at load of 0.5 N. The load was applied linearly up to the maximum load in 60 s followed by unloading in 60 s. The hardness and elastic modulus were also calculated by the Oliver and Pharr method [16]. A total of 10 indentations with the same load were carried out on each sample. Fracture toughness (Kc) of the coatings was measured on the well-polished surface of the coatings by the indentation technique using the same micro-indentation tester with a Vickers indenter. The loads with 5 N and 49 N were used to generate the cracks in the coatings for the calculation of the fracture toughness using the following equation [17]
2=5 E P Kc ¼ v 1=2 H al
273
consisting of oxygen and argon with a volume ratio of 22:78 because the initial oxidation temperature of C/C composites was about 370 °C [18]. Weight changes of the coated C/C samples related to their temperature were recorded with thermogravimetric mode. Isothermal oxidation test of the coated C/C samples was carried out at 1500 °C in the corundum tube furnace with natural convection of air to investigate their isothermal oxidation behavior. After the furnace was heated up to 1500 °C, the coated C/C samples were put directly into the furnace, whereafter they were taken out at the designated time and cooled at room temperature in air for weighing. After that they were put into the furnace again for the next isothermal oxidation test. An electronic balance with a sensitivity of ±0.1 mg was used to calculate the weight change (W) of the samples by the following equation:
W¼
m1 m0 S
ð3Þ
where m0 and m1 are the weight of the coated samples before and after oxidation, respectively, S is the surface area of the coated samples. The morphology and microstructure of the samples were analyzed by a scanning electron microscopy (SEM, Quanta 600FEG, FEI, USA), equipped with energy dispersive X-ray spectroscopy (EDS). 3. Results and discussion Fig. 1(a) is a typical SEM surface image showing the morphology of the porous SiC nanoribbon layer on C/C composites. It is obvious that the porous layer exhibits an open and porous structure consisting of a large quantity of wire-like and curvature nanostructures with typical lengths in the range of several tens to several hundreds of micrometers; some of them even have lengths on the order of millimeters. From high-magnification SEM images
ð1Þ
where P is the applied load, E is the elastic modulus, H is the hardness, a is the half diagonal of the indents, l is the radial crack length, and v is the constant. In this study, v = 0.0122 because l/a ratios were between 1 and 2.5. A total of 5 indentations were performed on each sample. Tensile test was carried out using an Instron universal test machine (CMT5304-30 kN, Instrument Co., Ltd., China) to estimate the interfacial bonding strength between the coating and C/C substrate (ASTM standard C633). The cylin-drical aluminum rods with dimensions of U25 mm 25 mm were used as matching parts. The coated C/C samples were bonded with the matching parts at both ends by epoxy resins. After being solidified for 72 h at room temperature, they were tested by the universal test machine. The maximum force of each sample was recorded. The bonding strength (r) is calculated using the following equation:
r¼
F S
ð2Þ
where F is the maximum force and S is the cross-sectional area of the samples. A total of 3 samples were measured for each group. Thermogravimetric analysis (TGA) of the coated C/C samples was performed using a Metter Toledo Star TGA/SDTA 851 thermal analyzer to investigate their oxidation behavior between room temperature to 1500 °C. The coated C/C samples were heated from 300 °C to 1500 °C with a heating rate of 20 °C/min in simulated air
Fig. 1. (a) SEM surface image of the porous SiC nanoribbon layer on C/C composites; (b) high-magnification image of (a).
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Fig. 2. (a) SEM surface image of the coating with SiC nanoribbons; (b) cross-section backscattering electron image of the coating with SiC nanoribbons.
(Fig. 1(b)), a high width-to-thickness ratio of the nanoribbons can be clearly observed. The nanoribbons have smooth surfaces along their length and width and most of them have uniform rectangular cross sections with depths of 10–100 nm and widths of 100– 900 nm. Fig. 2(a) shows that SEM surface image of the coating with SiC nanoribbons. Obviously, the coating exhibited a dense surface morphology, on which the free Si was well distributed around SiC particles and a small quantity of Al2O3 was observed. Cross-section backscattering electron image (Fig. 2(b)) displays that the coating had a dense microstructure with the major SiC phase and Si phase, in which Si phase was well filled into the gaps among SiC phase and few holes and penetrating cracks were observed. SEM images shown in Fig. 3 display the morphology of the fracture surface of SiC nanoribbon-reinforced SiC–Si coating. From Fig. 3(a), the morphology of the fracture surface of the coating shows clearly a stepwise feature, indicating the crack occurred an obvious deflection during the coating failure. Of particular interest is that there are numerous SiC nanoribbons on the fracture surface. Most of SiC nanoribbons are observed to debond from the matrix, as exemplified in Fig. 3(b). These nanoribbons are partially exposed to the crack plane, and the exposed area of such nanoribbons lies parallel to the crack plane. Therefore, such nanoribbons induce the deflection of the crack during crack propagation, leading to the formation of the stepwise fracture surface. In addition, note that some SiC nanoribbons exhibit an obvious pullout feature on the fracture surface, which can absorb vast amounts of fracture energy due to the presence of the good connectivity between the nanoribbons and coating matrix. Fig. 3(c) is a typical SEM image showing bridging feature of SiC nanoribbon within the crack that resists the crack widening (red1 arrow in Fig. 3(c)), suggesting a 1 For interpretation of color in Figs. 3–5, 7 and 8, the reader is referred to the web version of this article.
potential crack bridge toughening mechanism. Moreover, Fig. 3(c) also shows that the crack propagates to SiC nanoribbon and then deflects to another direction (blue arrow in Fig. 3(c)), confirming that there is indeed the toughening mechanism of crack deflection in this coating. As shown in Fig. 3(d), the cracks showed markedly deflection near the nanoribbons when propagating near the nanoribbons, and then propagated along the nanoribbon boundary. Due to the good connectivity between the matrix and the nanoribbons, the propagation of crack would be resisted and turned to another direction when propagating along the nanoribbon boundary to a certain degree (red arrows in Fig. 3(d)). This process results in an increase in the propagation path and resistance of the crack, which can cause the coating to be toughened. The elastic modulus and hardness values of the coatings obtained from the micro-indentation tests are shown in Table 1. It can be found that the hardness and elastic modulus of the coating increase by 27% and 39% with the incorporation of SiC nanoribbons, respectively. The fracture toughness of the coatings was also measured by micro-indentation technology. First, the fracture toughness measurement at a low load was evaluated on the cross-section of the coatings to avoid substrate effect. Fig. 4(a) shows SEM image of the indentation generated at 5 N load on the cross-section of the coating without SiC nanoribbons. The radial cracks which propagated from the indentation corner, where the highest stresses occurred, were observed. According to the above results, the fracture toughness of the coatings can be calculated using Eq. (1). The result is presented in Table 1. The coating without SiC nanoribbons shows an average fracture toughness of 7.28 ± 0.85 MPam1/2. SEM image of the indentation generated at 5 N load on the cross-section of the coating with SiC nanoribbons was shown in the left-bottom of Fig. 4(b). An indentation with the diagonal length of 20 lm was observed, but no radial crack was found around the indentation. Considering the thickness of the coating has only 85 lm, the fracture toughness measurement at a high load was evaluated on the surface of the coatings. SEM image of the indentation generated at 49 N load on the surface of the coating with SiC nanoribbons is presented in Fig. 4(b). It was found that the diagonal length of the indentation was about 100 lm, which was larger than the thickness of the coating. It should be noted that the radial cracks propagated from the indentation corner were observed, and the fracture toughness of the coatings was calculated using the Eq. (1), which was presented in Table 1. The coating with SiC nanoribbons shows an average fracture toughness of 14.42 ± 0.97 MPam1/2. An increase in 98% in relative fracture toughness of the coating was observed with the incorporation of SiC nanoribbons. The reinforcement effectiveness of SiC nanoribbons is better than the reported results about SiC nanowires in our previous work [6,8,9]. Detailed SEM investigations inside the indentation cracks reveal the following factors responsible for the improvement in fracture toughness of the coatings. From Fig. 4(b), it is interesting to find that the radial cracks exhibit the zigzag propagation path in the coating. High-magnification SEM image (Fig. 4(c)) shows that the crack exhibits obvious deflection feature when it encounters SiC nanoribbons (red arrow in Fig. 4(c)) during its propagation, which results in the formation of the zigzag propagation path and finally contributes to an increase in the toughness of the coating. Furthermore, Fig. 4(d) shows the high-magnification image within the crack which clearly shows numerous nanoribbon pullout and bridged features within the crack that resist the crack widening (red arrow in Fig. 4(d)) and hence contribute toward the toughening of the coating. In addition to the fracture toughness of the coatings, the interfacial bonding between the coatings and C/C composites is also regarded as the key factor to determine the oxidation resistant capacity of the coated C/C composites. The good interfacial bonding can alleviate thermal stress between the coatings and C/C composites, which ultimately affects the oxidation resistant
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Fig. 3. SEM fracture surface images of SiC nanoribbon-reinforced SiC–Si coating: (a) a representative SEM image of numerous debonding and pullout nanoribbons; (b) a representative SEM image of numerous debonding nanoribbons (magnified image of (a)); (c) a representative SEM image of crack bridging by the nanoribbon and crack deflection near the nanoribbon; (d) a representative SEM image of crack deflection near the nanoribbon.
Table 1 Micro-mechanical properties of the coatings. Samples
Hardness (GPa)
Elastic modulus (GPa)
Fracture toughness (MPam1/2)
Bonding strength (MPa)
Coating Nanoribbon-coating
24.53 ± 0.64 31.04 ± 0.46
258.89 ± 7.39 358.58 ± 11.37
7.28 ± 0.85 14.42 ± 0.97
21.53 ± 1.61 42.86 ± 2.54
Fig. 4. (a) SEM image of 5 N Vickers indentations on the cross-section of the coating without SiC nanoribbons; (b) SEM image of 49 N Vickers indentations on the surface of the coating with SiC nanoribbons (the inset is SEM image of 5 N Vickers indentations on the cross-section of the coating with SiC nanoribbons); (c) the representative crack deflection feature near the nanoribbons (enlarged image of the area marked as ‘‘A’’ in (b)); (d) the representative nanoribbon pullout and bridged features within the crack (enlarged image of the area marked as ‘‘B’’ in (b)).
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capacity of the coated C/C composites [19]. Table 1 presents a comparative study of the interfacial bonding strength between C/C composites and coatings. After incorporating SiC nanoribbons, the interfacial bonding strength between the coating and C/C composites represents a 99% increase (from 21.53 ± 1.61 to 42.86 ± 2.54 MPa). SEM fracture surface image of the interfacial area between the coating and C/C substrate is displayed in Fig. 5 (a). Numerous pullout and debonding SiC nanoribbons were observed on the fracture surface. More importantly, some SiC nanoribbon bridging features were detected between the coating and C/C substrate (red arrows in Fig. 5(a)), which implied the presence of interface anchoring effect of SiC nanoribbon between the coating and C/C substrate. A schematic diagram of in situ grown SiC nanoribbons on the interfacial area between the coating and C/C substrate is shown in Fig. 5(b). It is evinced here that some in situ grown SiC nanoribbons in the pores of C/C substrate established the anchoring effect between SiC nanoribbons and C/C substrate, as shown in Fig. 5(c) (high-magnification of Fig. 5(b)). These SiC nanoribbons may attach themselves strongly to the surrounding fibers. During SiC nanoribbon pullout, the relative sliding can occur between SiC nanoribbons and the surrounding fibers, which will resist the pullout of SiC nanoribbons, resulting in an improvement in the interfacial bonding between the coating and C/C substrate. Fig. 6(a) shows that the weight change curve of the coated C/C samples with SiC nanoribbons during TGA test from 300 °C to
1500 °C. It can be seen that the oxidation of the samples was a continuous weight gain process and the final weight gain was 61.84 gm2 during TGA test, which suggested that the as-prepared coating had good oxidation protective ability for C/C composites between room temperature and 1500 °C. The season for this is because the incorporated SiC nanoribbons effectively suppressed the cracking of the coating by via various toughening mechanisms including nanoribbon pullout, microcrack bridging by nanoribbon and microcrack deflection, avoiding the formation of the penetrating cracks in the coating, as shown in Fig. 2(b). The weight gain of the samples was believed to due to the oxidation of the coating, which could be divided into two stages, marked as A and B, respectively. During the stage A (300–1400 °C), the weight gain of the samples was a slow process in that the oxidation rate of SiC and Si phases in the coating was slow as the temperature was lower than 1400 °C. As the temperature was over 1400 °C, the oxidation rate of SiC and Si phases in the coating would be accelerated. Therefore, the weight gain of the samples exhibited a rapid process during the stage B (1400–1500 °C). Fig. 6(b) displays the weight change of the coated C/C samples with SiC nanoribbons during isothermal oxidation test at 1500 °C in air. Clearly, the oxidation of the samples was also a continuous weight gain process during the isothermal oxidation test, which is quite consistent with a parabolic law. As shown in Fig. 6(c), square of the weight gain as a function of oxidation time, the linear relationship confirms that the oxidation follows parabolic behavior, suggesting that the
Fig. 5. (a) SEM fracture surface image of the interfacial area between the coating and C/C substrate; (b) and (c) schematic description of the anchoring mechanism of in situ grown SiC nanoribbons on the interfacial area between the coating and C/C substrate.
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277
Weight gain (g•m-2)
70 60
(a)
50 40 30 20
B
A
10 0 300
600
900
1200
1500
Weight gain (g•m-2)
Temperature/
40
(b)
30 20 10 0 0
20
40
60
80
100
Weight gain (g 2 •m -4 )
Oxidation time (h) 1600 1400 1200 1000 800 600 400 200 0
(c)
Fig. 7. (a) A representative SEM image of nanoribbon pullout within the crack on the surface of the coating with SiC nanoribbons after isothermal oxidation test; (b) SEM cross-section image of coated C/C samples with SiC nanoribbons after isothermal oxidation test (the inset is EDS analysis of spot A in (b)).
0
20
40
60
80
100
Oxidation time (h) Fig. 6. (a) Weight change curve of the coated C/C samples with SiC nanoribbons during TGA test with the increasing temperature from 300 °C to 1500 °C at the rate of 20 °C/min in air (the total oxidation time: 1 h); (b) weight change of the coated C/C samples with SiC nanoribbons as a function of oxidation time during isothermal oxidation test at 1500 °C in air; (c) square of the weight gain of the coated C/C samples with SiC nanoribbons as a function of oxidation time during isothermal oxidation test at 1500 °C in air.
oxidation is controlled by a diffusion process. The parabolic oxidation rate constants, kp, was calculated to be 13.52 gm2h1, less than that of the reported pure SiC ceramic [20], from the slope of the lines according to the following equation
W 2 ¼ kp t
ð4Þ
where W is the weight change of the samples, t the oxidation time. Meanwhile, the weight gain of the samples was 38.0 gm2 after isothermal oxidation at 1500 °C in air for 96 h, which is also less than that that of the samples after TGA test. These suggested that the weight gain of the samples was a combination of the weight gain resulting from the oxidation of the coating and the weight loss resulting from the oxidation of C/C substrate during isothermal oxidation. Obviously, the oxidation of C/C substrate was not dominant but slight and the oxidation of the coating was dominant during isothermal oxidation. Therefore, it can be inferred that the as-obtained coating has the potential to protect C/C composites against oxidation in a large temperature range. Fig. 7(a) shows SEM surface image of the coated C/C samples with SiC nanoribbons after isothermal oxidation test. A smooth SiO2 glass layer was formed on the coating surface (blue arrows
in Fig. 7(a)). Meanwhile, some cracks were observed, in which there are some pullout SiC nanoribbons (red arrow in Fig. 7(a)) and SiO2 glass (blue arrows in Fig. 7(a)), implying that they played a key role in suppressing the cracking of the crack and sealing the crack itself. Fig. 7(b) shows SEM cross-section image of the coated C/C samples with SiC nanoribbons after isothermal oxidation test, on which the following features can be found. First, a SiO2 glass layer was observed on the coating surface, in which there are some perfect bubbles and broken bubbles (red arrow in Fig. 7(b)) due to the oxidation of the coating and C/C substrate [9]. Second, the cracks were detected, which had penetrated the coating and extended into C/C substrate, resulting in a slight oxidized zone without void in C/C substrate. Of particular interest is that the penetrating cracks were not open but sealed with some materials. EDS analysis (Fig. 7(b)) shows that these materials involve Si, O, Al and Pt elements. The presence of Si, O and Al elements confirms that these materials are the generated SiO2 glass [15], indicating that the penetrating cracks can be sealed by the generated SiO2 glass at elevated temperatures. The presence of Pt element can be regarded as an impurity from the spray-gold for SEM analysis. Therefore, it can be inferred that the oxidation of C/C was controlled by the diffusion rate of O2 in the generated glass layer within the penetrating cracks in the coating at elevated temperatures. To explain the oxidation mechanisms of the coated C/C samples with SiC nanoribbons during isothermal oxidation, a schematic diagram is presented in Fig. 8. First, the samples would undergo thermal cycles for 6 times between 1500 °C and room temperature during the isothermal oxidation test for weighing. During this process, the large thermal stress would induce the cracking of the coating due to the mismatch of the coefficient of thermal expansion between the coating and C/C, however, the cracking of the coating
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gases would be accumulated together with the gases generated by the oxidation of the coating to form the bubbles in the generated glass layer, as shown in Fig. 8(b). The bubbles would break as the pressure of these gases exceeded the surface tension of the bubbles. This was also the season why the broken bubbles were found preferentially on the coating surface along the penetrating cracks (red arrow in Fig. 7(b)). It should be noted that the oxidation of C/C substrate is slight during isothermal oxidation in that it is controlled by the diffusion rate of O2 in the generated glass layer. 4. Conclusions In this study, a novel SiC nanoribbon-reinforced SiC–Si ceramic coating had been prepared on C/C composites by a two-step technique consisting of CVD and pack cementation. The as-prepared coating had a dense microstructure, in which SiC nanoribbons were well dispersed. The incorporated SiC nanoribbons significantly improved the hardness, elastic modulus and fracture toughness of the coating, which was primarily attributable to the nanoribbon pullout, nanoribbon bridging, crack deflection. Meantime, the incorporated SiC nanoribbons also greatly improved the interfacial bonding of the coating-substrate, which was mainly resulted from the anchoring effect of SiC nanoribbons at the coating-substrate interface. The as-received coating possessed excellent oxidation protective capacity for C/C composites in a large temperature range. This study provides a promising way to improve the oxidation protective capacity of ceramic coatings for carbon materials by the reinforcement effects of SiC nanoribbons. Acknowledgements This work has been supported by the National Natural Science Foundation of China (Grant Nos. 51221001, 51272212, and 51275417), the ‘‘111’’ Project (Grant No. B08040), the State Scholarship Fund of China Scholarship Council, the Doctorate Foundation of Northwestern Polytechnical University and Excellent Doctorate Foundation of Northwestern Polytechnical University. References
Fig. 8. Schematic description of possible oxidation mechanisms of the coated C/C samples with SiC nanoribbons during isothermal oxidation.
could be effectively suppressed by via various toughening mechanisms of SiC nanoribbons including nanoribbon pullout, microcrack bridging by nanoribbon and microcrack deflection. Thus, the size and number of the cracks in the coating were decreased to a large extent. At the same time, the generated SiO2 glass possessed the good fluidity at elevated temperatures, which could effectively seal the generated penetrating cracks in the coating, as shown in Fig. 8(a). Thus, the as-obtained coating has a good oxidation protective ability for C/C composites. While the presence of Al3+ could damage the network structure of the SiO2 and result in the increase of oxygen diffusion in the generated SiO2 glass [9]. For one thing some oxygen diffused easily to the surface of the coating through the glass layer (indicated by red arrows), continued to react with the coating and generated some gases such as CO, CO2 and SiO (indicated by blue arrows in Fig. 8(b)) [9]. For another, some oxygen also diffused easily to C/C substrate through the generated SiO2 glass within the penetrating cracks in the coating (indicated by red arrows), leading to the oxidation of C/C substrate and releasing CO2 and CO gases (indicated by blue arrows in Fig. 8(b)) [9]. These
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