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ZrC-SiC-based composites fabricated by an improved reactive melt infiltration
Journal of the European Ceramic Society 39 (2019) 4617–4624
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Original article
Ablation behavior of Cf/ZrC-SiC-based composites fabricated by an improved reactive melt infiltration
T
⁎
Bo-Wen Chena,c,d, De-Wei Nia,c, , Jing-Xiao Wanga,c, You-Lin Jianga,c,d, Qi Dinga,c, Le Gaoa,c, ⁎ Xiang-Yu Zhanga,c, Yu-Sheng Dinga,c, Shao-Ming Donga,b, a
State Key Laboratory of High Performance Ceramics and Superfine Microstructure, Shanghai Institute of Ceramics, Chinese Academy of Sciences, Shanghai, China Center of Materials Science and Optoelectronics Engineering, University of Chinese Academy of Sciences, Beijing 100049, China c Structural Ceramics and Composites Engineering Research Center, Shanghai Institute of Ceramics, Chinese Academy of Sciences, Shanghai, China d University of Chinese Academy of Sciences, Beijing, China b
A R T I C LE I N FO
A B S T R A C T
Keywords: Composites UHTCs Reactive melt infiltration Ablation resistance
In this work, a highly dense Cf/ZrC-SiC-based composite is fabricated by an improved reactive melt infiltration (RMI). The ablation resistance of the composite is studied by air plasma test. The RMI-Cf/ZrC-SiC possesses a low porosity (3.49%) and high thermal conductivity. The dense microstructure can effectively retard oxygen from diffusing into the interior composite. Meanwhile, the high thermal conductivity makes the composite transfer heat timely during ablation, which reduces the heat accumulation on the ablation surface and weakens the thermal damage to the composite. Consequently, the as-fabricated composite exhibits an excellent ablation resistance. Compared to conventional PIP-Cf/ZrC-SiC composite, the linear and mass recession rates of the RMICf/ZrC-SiC decline by 98.07% and 39.02% at a heat flux of 4.02 MW/m2. Also, a continuous SiO2-ZrO2 layer forms on the sample surface, which isolates the sample surface from the plasma flame and protect the composites from further oxidation and ablation.
1. Introduction Continuous carbon fibers reinforced ultra-high temperature ceramic matrix composites (Cf/UHTCs), such as Cf/ZrC [1], Cf/ZrC-SiC [2,3], Cf/ZrC-SiC-ZrB2 [4], Cf/HfB2-SiC [5], etc., possess excellent mechanical properties and ablation resistance at ultra-high temperature, and fundamentally overcome the brittleness and poor thermal shock resistance of bulk UHTCs. These outstanding properties of the Cf/UHTCs make them key candidates for wide applications as high-temperature structure materials and thermal protection system of the hypersonic vehicle [6]. Particularly, Zr-based Cf/UHTCs, such as Cf/ZrC-SiC and Cf/ZrB2SiC, have received a lot of attention in the past few years due to their lower cost and lower density than Hf-based Cf/UHTCs [7,8]. Up to now, a lot of works have been devoted to studying the fabrication and the properties of Cf/UHTCs. Chemical vapor infiltration (CVI) [9,10], precursor infiltration and pyrolysis (PIP) [11–13], slurry infiltration (SI) [14,15] and reactive melt infiltration (RMI) [16,17], etc. are the most common methods used in fabrication of Cf/UHTCs. Compared with CVI, PIP and SI methods, RMI is more efficient and allows reaching a high densification level in the composites [18].
However, the composites fabricated by conventional RMI method generally appear fibers/interphase degradation and large-sized metal residues due to the limitation of preparation processing [19]. As a result, the reliability and service life of the composites are seriously restricted. Based on ZrC interphase-protective-structure and sol-gel pore structures tailoring, highly dense Cf/ZrC-SiC-based composites without evident fiber/interphase degradation were fabricated in our previous work. The as-fabricated Cf/ZrC-SiC composites show much higher mechanical properties than the composites prepared by conventional RMI method [20]. Ablation resistance at high temperature is one of the most critical properties required for the applications of Cf/UHTCs [4,7,12,21–23]. Some studies well considered the ablation and oxidation behavior of ZrC-SiC monoliths [24,25], however, the ablation behavior of Cf/ UHTCs has not been well understood until now, due to the extremely complicated physical-chemical processes occurring during ablation. In this work, the ablation behavior of the Cf/ZrC-SiC composite fabricated by the improved reactive melt infiltration method was studied and compared with the conventional PIP-Cf/ZrC-SiC. Particularly, the internal relation and the mechanisms among the microstructure,
⁎ Corresponding authors at: State Key Laboratory of High Performance Ceramics and Superfine Microstructure, Shanghai Institute of Ceramics, Chinese Academy of Sciences, Shanghai, China. E-mail addresses: [email protected] (D.-W. Ni), [email protected] (S.-M. Dong).
https://doi.org/10.1016/j.jeurceramsoc.2019.07.007 Received 12 April 2019; Received in revised form 28 June 2019; Accepted 3 July 2019 Available online 04 July 2019 0955-2219/ © 2019 Elsevier Ltd. All rights reserved.
Journal of the European Ceramic Society 39 (2019) 4617–4624
B.-W. Chen, et al.
Fig. 1. SEM images of RMI-Cf/ZrC-SiC (a), (c), (e), (f) and PIP-Cf/ZrC-SiC (b), (d), (g) composites.
ablated surface was measured by a bichromatic pyrometer (RAYMR1SCSF, Raytek, USA). The ablation recession rates were calculated based on the mass and thickness variations before and after ablation test according to the following equations:
thermophysical characteristics and ablation-resistant properties of the composites are revealed. 2. Experimental procedure
MA =
2.1. Materials and processing
m 0 − mt t
(1)
where MA is the mass recession rate (mg/s), m0 and mt are the mass (mg) of composites before and after ablation, t is ablation time (s).
Three-dimensional (3D) carbon fibers (PAN-based T700, Toray Industry) was needled as the reinforcement of the composites. The obtained carbon fiber fabric has a density of 0.52 g/cm3 and fiber volume fraction of about 32 vol%, which was first deposited with (PyCSiC)3 interphase by CVI process. And then, a Cf/ZrC-C nano-porous preform was prepared by impregnation of ZrC-C sol, followed by heat treatment. Finally, the Cf/ZrC-SiC composite was obtained by RMI of Si melt into the nano-porous Cf/ZrC-C preform at 1500 ℃ in the vacuum. Details of the fabrication process can be found in our previous work [20]. The PIP-Cf/ZrC-SiC composite was fabricated by conventional precursor infiltration and pyrolysis method for comparison.
RA =
l 0 − lt t
(2)
where RA is linear recession rate (μm/s), l0 and lt are thickness (μm) of composites before and after ablation. Specific heat and thermal diffusivity were measured by laser flash thermal conductivity meter (NETZSCH LFA467 HT, Germany). The size of the sample used for thermal diffusivity measurement is Φ10 mm × 2 mm. Thermal conductivity was calculated according to the equation below:
λ = αCp ρ
2.2. Test and characterization
(3) −1
−1
where λ and α are thermal conductivity (W⋅ m ⋅ K ) and thermal diffusivity (mm2⋅s−1), Cp and ρ are specific heat (J⋅ g−1⋅ K−1) and density (g⋅ cm-3) of the composites. The densities and open porosities of the composites were measured by the Archimedes’ method using deionized water as immersing medium. Phase compositions of the composites before and after
The as-fabricated composites were processed into a rectangular shape (˜60 mm × 30 mm). The plasma torch (Plasmer Jet, A-2000, Sulzer Metco, Switzerland) was used to test the ablation resistance of the specimens. The heat fluxes were 3.01 and 4.02 MW/m2, and the testing time was 150 s and 60 s respectively. The temperature of the 4618
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ablation test were determined by X-ray diffraction (XRD, Ultima IV, Rigaku Corporation, Kyoto, Japan) using Cu-Kα radiation. Microstructural features of the samples were characterized by scanning electron microscopy (SEM, S4800, Tokyo, Japan) combined with energy dispersive spectroscopy (EDS, Oxford Instruments, UK). Based on the different grayscale of ZrC and SiC in SEM image, the volume fractions of ZrC and SiC in composites were determined by ImageJ® freeware (https://imagej.nih.gov/ij/) through image analysis. To reduce the error of image analysis, at least 10 images were taken in random regions of the samples and used for statistical analysis. 3. Result and discussion 3.1. Microstructure and compositions The as-fabricated RMI-Cf/ZrC-SiC composite has a density of 2.64 g/ cm3 and a porosity of 3.49%, whereas the density and open porosity of the PIP-Cf/ZrC-SiC composite are 2.18 g/cm3 and 13.53% respectively. Fig. 1 illustrates the polished cross-sections of the composites. In agreement with the density measurement, RMI-Cf/ZrC-SiC shows a highly dense microstructure, with almost no obvious pores. Three different phases present in the matrix of the composites: bulk ZrC, white needle-shaped ZrSi2 and continuous dark grey SiC, as indicated in Fig. 1c according to EDS analysis. As reported before, the needle-shaped ZrSi2 is formed via the reaction between ZrC and Si during the cooling stage of RMI [20]. It can also be observed that fibers and interphase are well protected in this RMI-Cf/ZrC-SiC composite, without obvious fibers/ interphase erosion (Fig. 1e). In comparison, the PIP-Cf/ZrC-SiC composite presents a completely different microstructure. Quite a few small pores can be observed on the polished cross-section of the PIP-Cf/ ZrC-SiC composite, where the ZrC and SiC show a uniform distribution with fine particle size (Fig.1d and g). The volume fractions of ZrC and SiC in the two composites are determined based on image analysis and illustrated in Table 1. SiC is the main phase component in both composites. But the SiC content in the PIP-Cf/ZrC-SiC is a little bit less than that in the RMI-Cf/ZrC-SiC composite. Fig. 2 displays the XRD patterns of the as-fabricated composites, where the results are consistent with the SEM and EDS analysis. The RMI-Cf/ZrC-SiC composite is mainly composed of ZrC, SiC and ZrSi2, while the PIP-Cf/ZrC-SiC composite is composed of ZrC and SiC phases. In addition, the wide diffraction peak, located at around 26°, corresponds to carbon fibers of the composites.
Fig. 2. XRD patterns of composites before ablation.
the measured density, thermal diffusivity and specific heat, the calculated thermal conductivity of RMI-Cf/ZrC-SiC composite decreases gradually from 21.26 W⋅ m-1⋅ K-1 at room temperature to 15.43 W⋅ m1 ⋅ K-1 at 1200 °C. The thermal conductivity of RMI-Cf/ZrC-SiC composite is quadruple higher than that of the PIP-Cf/ZrC-SiC composite (Fig. 3c). The high thermal conductivity of the RMI-Cf/ZrC-SiC composite can be attributed to its low porosity (3.49%) and continuous phase distribution. The relation between thermal conductivity and porosity can be expressed roughly by the following equation [26].
λ = λs (1 − p)
Thermal diffusivity, specific heat and thermal conductivity of the composites were measured in the temperature range of RT-1200 °C. The results are shown in Fig. 3. The thermal diffusion coefficient is a typical physical quantity that describes how fast the heat is transferred in materials. The thermal diffusion coefficient of the as-fabricated RMI-Cf/ ZrC-SiC composite is 12.06 mm2·s−1 at room temperature, which decreases to 4.20 mm2·s−1 at 1200 °C. Specific heat is used to evaluate the energy storage capacity of materials. As shown in Fig. 3b, the specific heat of the RMI-Cf/ZrC-SiC composite increases gradually with the increase in temperature and reaches 1.47 J⋅ g-1⋅ K-1 at 1200 °C. Based on
3.3. Ablation-resistant property and mechanisms Fig. 4 illustrates the macroscopic features of the composites after ablation under different heat fluxes. The ablated surface can be divided into central ablation area and transition area. This typical ablation morphology is caused by the synergic effects of heat, oxidation and mechanical scouring. There’s no distinct ablation damage on the RMICf/ZrC-SiC composite (Fig. 4a-b). As shown in Fig. 4a, only a small amount of oxide spatters out from the ablation center after ablation by the high heat flux (4.02 MW/m2). In comparison, the PIP-Cf/ZrC-SiC composite is damaged remarkably and a specific ablation center can be clearly observed (Fig. 4c-d).
Table 1 Parameters of the RMI-Cf/ZrC-SiC and the PIP- Cf/ZrC-SiC composites.
RMI-Cf/ZrCSiC PIP- Cf/ZrCSiC
Density (g/ cm3)
Open porosity (%)
Phase volume fraction (%) ZrC
ZrSi2
SiC
Cf
2.64
3.49
6.69
6.76
50.52
32.54
2.18
13.53
13.22
N/A
41.20
32.05
−1
where λ is the thermal conductivity (W⋅ m ⋅ K ), λs is the thermal conductivity (W⋅ m−1⋅ K−1) of fully dense material, and p is the porosity (%) of the material. Assuming the RMI-Cf/ZrC-SiC composite has the same porosity (13.53%) as the PIP-Cf/ZrC-SiC composite, the thermal conductivity with 13.53% of porosity in the RMI-Cf/ZrC-SiC composite is calculated to be 19.05 W⋅ m−1⋅ K−1 at room temperature and 13.83 W⋅ m−1⋅ K−1 at 1200 ℃. It can be found that the estimated thermal conductivity values are also higher than those of the PIP-Cf/ZrC-SiC composite. This indicates that in addition to its low porosity, other factors also contribute to the high thermal conductivity of the RMI-Cf/ZrC-SiC composite. Generally, heat is conducted mainly through phonon motion of the solid framework (crystal lattice vibration) in composites [27]. Because of the dispersed distribution characteristics of ZrC particles in the PIPCf/ZrC-SiC composite (Fig. 1g), the SiC phase is separated into small pieces, which remarkably increases the number of phase boundaries. Consequently, the abundant phonons scatter at phase boundaries and high porosity (13.53%) leads to the low thermal conductivity of the PIP-Cf/ZrC-SiC composite.
3.2. Thermophysical properties
Samples
(4) −1
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Fig. 3. Thermal diffusivity (a), specific heat (b) and thermal conductivity (c) of RMI-Cf/ZrC-SiC and PIP-Cf/ZrC-SiC composites.
ablation resistance of the composites. As described in Fig. 6a, reactions (Eq. (7) and (8)) are more thermodynamically favorable at a lower temperature. Therefore, a continuous SiO2-ZrO2 layer tends to form on the ablated surface at the starting stage of the ablation test. With the increase of temperature, the chemical process of Eq. (5) and (6) becomes more active, and a mass of gaseous products (SiO and CO) are generated. The evaporation of SiO and CO results in the formation of a lot of pores in the SiO2-ZrO2 layer. The difference in microstructure of these two composites also gives influence to their oxidation behavior. The bulk ZrC (Fig. 1f) in the RMICf/ZrC-SiC composite is oxidized into continuous ZrO2 which dissolves in glassy SiO2. The continuous ZrO2 acts as a net structure to prevent SiO2 collapse from flame scouring and reduces the evaporation of SiO2 (Fig. 7a). In comparison, the dispersed distribution of ZrC (Fig. 1g) in PIP-Cf/ZrC-SiC is oxidized into dispersed ZrO2 which may has less pinning effect for glassy SiO2. So, the glassy SiO2 is easily scoured by flame (Fig. 7b). Fig. 6b shows the XRD patterns of the composites at ablation center. m-ZrO2 and SiC are detected in both the RMI-Cf/ZrC-SiC and the PIP-Cf/ ZrC-SiC composites at different heat flux conditions. It can be found that the diffraction peaks intensity of ZrO2 is relatively lower for the composites ablated at a heat flux of 3.01 MW/m2 due to limited ZrO2 formed under this condition. However, SiO2 is not detected in Fig. 6b. After ablation, the surface temperature drops sharply and the SiO2based glassy phase commonly stays in an amorphous state under this condition. As a result, no diffraction or diffusion peaks of the glassy phase are detected in the XRD spectra. This phenomenon has also been well reported in literature [4,28–30]. The existence of SiC indicates that SiC matrix is not completely covered by the glassy oxide layer. Some cracks and pores formed on the oxide layer during ablation (Fig. 7a and b), leading to the exposure of SiC. As shown in Fig. 7c and d, it can be observed that most of the sample surface is well covered by glassy layer at the ablated center, particularly for the RMI- Cf/ZrC-SiC. In order to understand the ablation behavior carefully, further observation was performed on the ablation area of the composites ablated at a heat flux of 4.02 MW/m2 for 60 s. The results are shown in Fig. 7.
The calculated ablation recession rates are plotted in Fig. 5. The mass and linear ablation rates of both composites at 4.02 MW/m2 are much higher than those of the composites at 3.01 MW/m2, indicating the composites are easier ablated at high heat flux, even tested for a shorter period (Fig. 5). The mass and linear ablation rates of the RMICf/ZrC-SiC composite decline by 39.02% and 98.07% respectively at 4.02 MW/m2 compared with the PIP-Cf/ZrC-SiC composite. The mass and linear ablation rates of the RMI-Cf/ZrC-SiC composite at 3.01 MW/ m2 are 33.53% and 98.71% smaller than those of the PIP-Cf/ZrC-SiC composite (Fig. 5). Therefore, RMI-Cf/ZrC-SiC composite shows a much better ablation-resistant property at two tested heat flux conditions. The flame temperature at high heat flux is about 3000 °C, which is much higher than that of low heat flux (˜2300 °C). Thus, the composites ablated at 4.02 MW/m2 have a higher surface temperature compare with the case of 3.01 MW/m2 (Table 2). Moreover, due to the much higher thermal conductivity, the heat on the ablated surface of the RMICf/ZrC-SiC composite can be transmitted and dissipated timely during the ablation, which reduces the surface temperature. As shown in Table 2, the surface temperature of the RMI-Cf/ZrC-SiC composite is 130–210 °C lower than that of the PIP-Cf/ZrC-SiC composite at the same ablation condition, which weakens the thermal damage of the RMI-Cf/ ZrC-SiC composite. Generally, the main reactions of the Cf/ZrC-SiC composites during the high-temperature ablation include: C + 0.5O2 (g) = CO (g)
(5)
SiC + O2 (g) = SiO (g) + CO (g)
(6)
SiC + 1.5O2 (g) = SiO2 + CO (g)
(7)
ZrC + 1.5O2 (g) = ZrO2 + CO (g)
(8)
The changes of standard Gibbs free energy of the above reactions are calculated as a function of temperature and plotted in Fig. 6a. The formation of a continuous SiO2-ZrO2 layer on the ablated surface can protect the composites from further oxidation and ablation. However, the evaporation of gaseous products during ablation is harmful to the 4620
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Fig. 4. Macroscopic features of the composites after ablation at different heat fluxes: (a) RMI-Cf/ZrC-SiC ablated at 4.02 MW/m2 for 60 s; (b) RMI-Cf/ZrC-SiC ablated at 3.01 MW/m2 for 150 s; (c) PIP-Cf/ZrC-SiC ablated at 4.02 MW/m2 for 60 s; (d) PIP-Cf/ZrC-SiC ablated at 3.01 MW/m2 for 150 s.
composite. By contrast, the dispersed distribution of ZrC in PIP-Cf/ZrCSiC composite divides SiC phase into small pieces along with a mass of small pores (Fig. 1d and g), which provide oxygen diffusing paths and result in the fragmentized oxide layer (Fig.7d). Fig. 7 e–f show the microstructure of the transition area on the ablated surface. As the temperature of the transition area is much lower than that of the ablation center, both samples are covered by the continuous glassy oxide layer, without evident signs left by massive gaseous product. As discussed above, the ablation behavior can be affected by microstructure, thermophysical properties and the formed oxide layer, etc. of the composites. Considering their interactions, the different ablationresistant property and mechanism between the RMI-Cf/ZrC-SiC and the PIP-Cf/ZrC-SiC composites can be summarized as Fig. 8:
After ablation test, the ablation center of the RMI-Cf/ZrC-SiC composite is covered by a continuous oxide layer, as shown in Fig. 7a and c. The EDS analysis (Fig. 7g) reveals that the continuous oxide layer is composed of SiO2 glass with ZrO2 particles embedding in it. Because SiO2 playing as a bonding agent, oxygen cannot easily diffuse through the continuous SiO2-ZrO2 layer. The continuous SiO2-ZrO2 layer plays a significant role in improving the ablation-resistant property as it can separate the surface of the composites from the ablation atmosphere and protect the underlying composites from damage. In comparison, the surface temperature of the PIP-Cf/ZrC-SiC composite is as high as 2100 °C during ablation test at a heat flux of 4.02 MW/m2, which is 210 °C higher than that of the RMI-Cf/ZrC-SiC composite (Table 2). As a result, SiO and CO are generated heavily by reactions (Eq. (5) and (6)) and escaped quickly through the surficial oxide layer. Meanwhile, SiO2 cannot survive due to its serious evaporation at such high temperature. Consequently, the ablation center of the PIP-Cf/ZrC-SiC composite exhibits a loose powdery microstructure, as shown in Fig. 7b and d. As a result, the protective effect of the surficial oxide layer on the PIP-Cf/ ZrC-SiC composite is weakened at such condition. On the other side, under the surface oxide layer, a porous layer presents in both the RMICf/ZrC-SiC and the PIP-Cf/ZrC-SiC composites, where the pores form due to the recession of SiC and carbon fibers during ablation. The difference in ablation-resistant property between the two composites is also related to their microstructure and phase constituent. The dense and continuous microstructure of the RMI-Cf/ZrC-SiC composite can effectively retard oxygen from diffusing into the interior of the composite and further slow down the ablation recession of the
(i) At the beginning of ablation, the exterior ZrC and SiC are oxidized and a continuous SiO2-ZrO2 glassy layer forms and covers on the sample surface. (ii) With the ablation process going on and the surface temperature increasing, the production and recession of SiO2-ZrO2 layer reach a dynamic equilibrium state. Benefiting from the low porosity and high thermal conductivity of RMI-Cf/ZrC-SiC composite, the heat is timely delivered from the ablation surface. Ablation center temperature is kept below 1900 °C even though the flame temperature is as high as 3000 °C at a heat flux of 4.02 MW/m2. Consequently, the continuous SiO2-ZrO2 protective layer is retained on the composite surface during the ablation process, leading to an excellent
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Fig. 5. Linear recession rate and mass recession rate under different ablation condition.
structure. (iv) The loose powdery structure of the surficial oxide cannot provide effective protection for the PIP-Cf/ZrC-SiC composite. Consequently, the recession of inside matrix and carbon fibers is aggravated, and the damage is going inward.
Table 2 Parameters of the ablation test. Ablation condition Heat flux (MW/m2)
Time (s)
4.02
60 s
3.01
150 s
Samples
Surface temperature (℃)
RMI-Cf/ZrC-SiC PIP- Cf/ZrC-SiC RMI-Cf/ZrC-SiC PIP- Cf/ZrC-SiC
˜1890 ˜2100 ˜1700 ˜1830
4. Conclusion Cf/ZrC-SiC-based composite is fabricated by an improved reactive melt infiltration (RMI). The microstructure and ablation-resistant property of the composite are investigated comprehensively, and compared with conventional PIP-Cf/ZrC-SiC composite. The RMI-Cf/ ZrC-SiC composite has low porosity (3.49%) and high thermal conductivity (21.3 W⋅ m−1⋅ K−1 at room temperature, 15.4 W⋅ m−1⋅ K−1 at 1200 °C), which is beneficial for transferring the heat timely and reducing the surface temperature of the composite during ablation. As a
ablation-resistant property for the RMI-Cf/ZrC-SiC composite. (iii) For PIP-Cf/ZrC-SiC composite, its limited thermal conductivity causes a sharp increase of the surface temperature. SiO2 evaporated quickly along with a heavy generation of SiO and CO gas. As a result, ZrO2 is left on the surface and forms a loose powdery
Fig. 6. (a) Changes of standard Gibbs free energy of reaction (5)-(8) as a function of temperatures calculated by HSC 6.0 software [31,32]; (b) XRD patterns of the composites at ablation center. 4622
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Fig. 7. SEM images of ablated surfaces and cross-sections at a heat flux of 4.02 MW/m2 for 60 s. (a, c) ablation center of RMI-Cf/ZrC-SiC; (b, d) ablation center of PIPCf/ZrC-SiC; (e, f) transition area of ablated RMI-Cf/ZrC-SiC and PIP-Cf/ZrC-SiC; (g) EDS analysis of the spots marked in (a), (b).
ZrC-SiC composite. During ablation, a continuous glassy SiO2-ZrO2 layer can be formed and cover on the ablation surface to hinder the ingression of oxygen into the RMI-Cf/ZrC-SiC composite. The layer also isolates the underlying materials from the plasma flame.
result, the RMI-Cf/ZrC-SiC composite presents an excellent ablationresistant property at two tested heat fluxes. The linear and mass recession rates of the RMI-Cf/ZrC-SiC composite decline by 98.71% and 33.53% at a heat flux of 3.01 MW/m2, and they decrease by 98.07% and 39.02% at a heat flux of 4.02 MW/m2 compared with the PIP-Cf/
Fig. 8. Schematic diagram of the ablation mechanism for Cf/ZrC-SiC composites. 4623
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[15] A. Vinci, L. Zoli, D. Sciti, Influence of SiC content on the oxidation of carbon fibre reinforced ZrB2/SiC composites at 1500 and 1650 °C in air, J. Eur. Ceram. Soc. 38 (2018) 3767–3776. [16] X.W. Chen, S.M. Dong, Y.M. Kan, H.J. Zhou, J.B. Hu, D.K. Wang, 3D Cf/SiC-ZrCZrB2 composites fabricated via sol-gel process combined with reactive melt infiltration, J. Eur. Ceram. Soc. 36 (2016) 3607–3613. [17] X.W. Chen, S.M. Dong, Y.M. Kan, X.H. Jin, H.J. Zhou, D.W. Ni, D.K. Wang, Microstructure and mechanical properties of three dimensional Cf/SiC-ZrC-ZrB2 composites prepared by reactive melt infiltration method, J. Eur. Ceram. Soc. 36 (2016) 3969–3976. [18] Q. Zhong, X.Y. Zhang, S.M. Dong, J.S. Yang, J.B. Hu, L. Gao, P. He, H.J. Zhou, Z. Wang, Y.S. Ding, Reactive melt infiltrated Cf/SiC composites with robust matrix derived from novel engineered pyrolytic carbon structure, Ceram. Int. 43 (2017) 5832–5836. [19] D.K. Wang, S.M. Dong, H.J. Zhou, X.Y. Zhang, Y.S. Ding, G.X. Zhu, Fabrication and microstructure of 3D Cf/ZrC-SiC composites: through RMI method with ZrO2 powders as pore-making agent, Ceram. Int. 42 (2016) 6720–6727. [20] D.W. Ni, J.X. Wang, S.M. Dong, X.W. Chen, Y.M. Kan, H.J. Zhou, L. Gao, X.Y. Zhang, Fabrication and properties of Cf/ZrC-SiC-based composites by an improved reactive melt infiltration, J. Am. Ceram. Soc. 101 (2018) 3253–3258. [21] L. Xue, Z.A. Su, X. Yang, D. Huang, T. Yin, C. Liu, Q. Huang, Microstructure and ablation behavior of C/C-HfC composites prepared by precursor infiltration and pyrolysis, Corros. Sci. 94 (2015) 165–170. [22] Y. Ma, Q.G. Li, S.M. Dong, Z. Wang, G.P. Shi, H.J. Zhou, Z. Wang, P. He, Microstructures and ablation properties of 3D 4-directional Cf/ZrC-SiC composite in a plasma wind tunnel environment, Ceram. Int. 40 (2014) 11387–11392. [23] Z. Li, H. Li, S. Zhang, J. Wang, W. Li, F. Sun, Effect of reaction melt infiltration temperature on the ablation properties of 2D C/C-SiC-ZrC composites, Corros. Sci. 58 (2012) 12–19. [24] L. Zhao, D. Jia, X. Duan, Z. Yang, Y. Zhou, Oxidation of ZrC–30vol% SiC composite in air from low to ultrahigh temperature, J. Eur. Ceram. Soc. 32 (2012) 947–954. [25] D. Pizon, L. Charpentier, R. Lucas, S. Foucaud, A. Maître, M. Balat-Pichelin, Oxidation behavior of spark plasma sintered ZrC–SiC composites obtained from the polymer-derived ceramics route, Ceram. Int. 40 (2014) 5025–5031. [26] J. Francl, W.D. Kingery, Thermal Conductivity: IX, Experimental Investigation of Effect of Porosity on Thermal Conductivity, J. Am. Ceram. Soc. 37 (1954) 99–107. [27] R.Y. Luo, T. Liu, J.S. Li, H.B. Zhang, Z.J. Chen, G.L. Tian, Thermophysical properties of carbon/carbon composites and physical mechanism of thermal expansion and thermal conductivity, Carbon 42 (2004) 2887–2895. [28] J. Xie, Y. Jia, Z. Zhao, K. Li, G. Sun, H. Li, X. Su, A ZrC SiC/SiC multilayer antiablation coating for ZrC modified C/C composites, Vacuum 157 (2018) 324–331. [29] S. Wang, H. Li, M. Ren, Y. Zuo, M. Yang, J. Zhang, J. Sun, Microstructure and ablation mechanism of C/C-ZrC-SiC composites in a plasma flame, Ceram. Int. 43 (2017) 10661–10667. [30] L. Duan, X. Zhao, Y. Wang, Comparative ablation behaviors of C/SiC-HfC composites prepared by reactive melt infiltration and precursor infiltration and pyrolysis routes, Ceram. Int. 43 (2017) 16114–16120. [31] I. Barin, O. Knacke, O. Kubaschewski, Thermodynamic Properties of Inorganic Substances, Springer-Verlag, Berlin, Heidelberg, New York, 1977. [32] L.V. Gurvich, Reference books and data banks on the thermodynamic properties of inorganic substances, in: J.W. Hastie (Ed.), Materials Chemistry at High Temperatures: Characterization, Humana Press, Totowa, NJ, 1990, pp. 197–206.
Acknowledgments The financial support from “The National Key Research and Development Program of China” [No. 2017YFB0703200], “National Natural Science Foundation of China” [No. 51702341, 51872310], “The Taicang Science and Technology Project” [No. TC2017DYDS21] and “CAS Pioneer Hundred Talents Program” are greatly acknowledged. References [1] D. Wang, Y. Wang, J. Rao, J. Ouyang, Y. Zhou, G. Song, Influence of reactive melt infiltration parameters on microstructure and properties of low temperature derived Cf/ZrC composites, Mater. Sci. Eng. A 568 (2013) 25–32. [2] Q. Li, S. Dong, Z. Wang, G. Shi, Y. Ma, H. Zhou, Z. Wang, Fabrication and properties of 3D Cf/ZrC–SiC composites via in-situ reaction, Ceram. Int. 40 (2014) 2483–2488. [3] J. Jiang, S. Wang, W. Li, Z. Chen, Y. Zhu, Preparation of 3D Cf/ZrC-SiC composites by joint processes of PIP and RMI, Mater. Sci. Eng. A 607 (2014) 334–340. [4] X. Chen, Q. Feng, H. Zhou, S. Dong, J. Wang, Y. Cao, Y. Kan, D. Ni, Ablation behavior of three-dimensional Cf/SiC-ZrC-ZrB2 composites prepared by a joint process of sol-gel and reactive melt infiltration, Corros. Sci. 134 (2018) 49–56. [5] S. Guo, K. Naito, Y. Kagawa, Mechanical and physical behaviors of short pitchbased carbon fiber-reinforced HfB2-SiC matrix composites, Ceram. Int. 39 (2013) 1567–1574. [6] C. Yan, R. Liu, C. Zhang, Y. Cao, X. Long, Mechanical behaviour and microstructure of Cf/ZrC, Cf/SiC and Cf/ZrC-SiC composites, Adv. Appl. Ceram. 115 (2016) 391–395. [7] X.T. Shen, L. Liu, W. Li, K.Z. Li, Ablation behaviour of C/C-ZrC composites in a solid rocket motor environment, Ceram. Int. 41 (2015) 11793–11803. [8] L. Zou, N. Wali, J.M. Yang, N.P. Bansal, Microstructural development of a Cf/ZrC composite manufactured by reactive melt infiltration, J. Eur. Ceram. Soc. 30 (2010) 1527–1535. [9] Y. Long, A. Javed, Y. Zhao, Z.K. Chen, X. Xiong, P. Xiao, Fiber/matrix interfacial shear strength of C/C composites with PyC-TaC-PyC and PyC-SiC-TaC-PyC multiinterlayers, Ceram. Int. 39 (2013) 6489–6496. [10] Z.K. Chen, X. Xiong, Microstructure, mechanical properties and oxidation behavior of carbon fiber reinforced PyC/C-TaC/PyC layered-structure ceramic matrix composites prepared by chemical vapor infiltration, Mater. Chem. Phys. 141 (2013) 613–619. [11] H. Zhong, Z. Wang, H.J. Zhou, D.W. Ni, Y.M. Kan, Y.S. Ding, S.M. Dong, Properties and microstructure evolution of Cf/SiC composites fabricated by polymer impregnation and pyrolysis (PIP) with liquid polycarbosilane, Ceram. Int. 43 (2017) 7387–7392. [12] W. Tan, K. Li, H. Li, J. Zhang, C. Ni, A. Cao, C. Ma, Ablation behavior and mechanism of C/C-HfC-SiC composites, Vacuum 116 (2015) 124–129. [13] C. Yan, R. Liu, Y. Cao, C. Zhang, Fabrication and properties of PIP 3D Cf/ZrC-SiC composites, Mater. Sci. Eng. A 591 (2014) 105–110. [14] A. Vinci, L. Zoli, D. Sciti, J. Watts, G.E. Hilmas, W.G. Fahrenholtz, Mechanical behaviour of carbon fibre reinforced TaC/SiC and ZrC/SiC composites up to 2100°C, J. Eur. Ceram. Soc. 39 (2019) 780–787.
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Original article
Ablation behavior of Cf/ZrC-SiC-based composites fabricated by an improved reactive melt infiltration
T
⁎
Bo-Wen Chena,c,d, De-Wei Nia,c, , Jing-Xiao Wanga,c, You-Lin Jianga,c,d, Qi Dinga,c, Le Gaoa,c, ⁎ Xiang-Yu Zhanga,c, Yu-Sheng Dinga,c, Shao-Ming Donga,b, a
State Key Laboratory of High Performance Ceramics and Superfine Microstructure, Shanghai Institute of Ceramics, Chinese Academy of Sciences, Shanghai, China Center of Materials Science and Optoelectronics Engineering, University of Chinese Academy of Sciences, Beijing 100049, China c Structural Ceramics and Composites Engineering Research Center, Shanghai Institute of Ceramics, Chinese Academy of Sciences, Shanghai, China d University of Chinese Academy of Sciences, Beijing, China b
A R T I C LE I N FO
A B S T R A C T
Keywords: Composites UHTCs Reactive melt infiltration Ablation resistance
In this work, a highly dense Cf/ZrC-SiC-based composite is fabricated by an improved reactive melt infiltration (RMI). The ablation resistance of the composite is studied by air plasma test. The RMI-Cf/ZrC-SiC possesses a low porosity (3.49%) and high thermal conductivity. The dense microstructure can effectively retard oxygen from diffusing into the interior composite. Meanwhile, the high thermal conductivity makes the composite transfer heat timely during ablation, which reduces the heat accumulation on the ablation surface and weakens the thermal damage to the composite. Consequently, the as-fabricated composite exhibits an excellent ablation resistance. Compared to conventional PIP-Cf/ZrC-SiC composite, the linear and mass recession rates of the RMICf/ZrC-SiC decline by 98.07% and 39.02% at a heat flux of 4.02 MW/m2. Also, a continuous SiO2-ZrO2 layer forms on the sample surface, which isolates the sample surface from the plasma flame and protect the composites from further oxidation and ablation.
1. Introduction Continuous carbon fibers reinforced ultra-high temperature ceramic matrix composites (Cf/UHTCs), such as Cf/ZrC [1], Cf/ZrC-SiC [2,3], Cf/ZrC-SiC-ZrB2 [4], Cf/HfB2-SiC [5], etc., possess excellent mechanical properties and ablation resistance at ultra-high temperature, and fundamentally overcome the brittleness and poor thermal shock resistance of bulk UHTCs. These outstanding properties of the Cf/UHTCs make them key candidates for wide applications as high-temperature structure materials and thermal protection system of the hypersonic vehicle [6]. Particularly, Zr-based Cf/UHTCs, such as Cf/ZrC-SiC and Cf/ZrB2SiC, have received a lot of attention in the past few years due to their lower cost and lower density than Hf-based Cf/UHTCs [7,8]. Up to now, a lot of works have been devoted to studying the fabrication and the properties of Cf/UHTCs. Chemical vapor infiltration (CVI) [9,10], precursor infiltration and pyrolysis (PIP) [11–13], slurry infiltration (SI) [14,15] and reactive melt infiltration (RMI) [16,17], etc. are the most common methods used in fabrication of Cf/UHTCs. Compared with CVI, PIP and SI methods, RMI is more efficient and allows reaching a high densification level in the composites [18].
However, the composites fabricated by conventional RMI method generally appear fibers/interphase degradation and large-sized metal residues due to the limitation of preparation processing [19]. As a result, the reliability and service life of the composites are seriously restricted. Based on ZrC interphase-protective-structure and sol-gel pore structures tailoring, highly dense Cf/ZrC-SiC-based composites without evident fiber/interphase degradation were fabricated in our previous work. The as-fabricated Cf/ZrC-SiC composites show much higher mechanical properties than the composites prepared by conventional RMI method [20]. Ablation resistance at high temperature is one of the most critical properties required for the applications of Cf/UHTCs [4,7,12,21–23]. Some studies well considered the ablation and oxidation behavior of ZrC-SiC monoliths [24,25], however, the ablation behavior of Cf/ UHTCs has not been well understood until now, due to the extremely complicated physical-chemical processes occurring during ablation. In this work, the ablation behavior of the Cf/ZrC-SiC composite fabricated by the improved reactive melt infiltration method was studied and compared with the conventional PIP-Cf/ZrC-SiC. Particularly, the internal relation and the mechanisms among the microstructure,
⁎ Corresponding authors at: State Key Laboratory of High Performance Ceramics and Superfine Microstructure, Shanghai Institute of Ceramics, Chinese Academy of Sciences, Shanghai, China. E-mail addresses: [email protected] (D.-W. Ni), [email protected] (S.-M. Dong).
https://doi.org/10.1016/j.jeurceramsoc.2019.07.007 Received 12 April 2019; Received in revised form 28 June 2019; Accepted 3 July 2019 Available online 04 July 2019 0955-2219/ © 2019 Elsevier Ltd. All rights reserved.
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Fig. 1. SEM images of RMI-Cf/ZrC-SiC (a), (c), (e), (f) and PIP-Cf/ZrC-SiC (b), (d), (g) composites.
ablated surface was measured by a bichromatic pyrometer (RAYMR1SCSF, Raytek, USA). The ablation recession rates were calculated based on the mass and thickness variations before and after ablation test according to the following equations:
thermophysical characteristics and ablation-resistant properties of the composites are revealed. 2. Experimental procedure
MA =
2.1. Materials and processing
m 0 − mt t
(1)
where MA is the mass recession rate (mg/s), m0 and mt are the mass (mg) of composites before and after ablation, t is ablation time (s).
Three-dimensional (3D) carbon fibers (PAN-based T700, Toray Industry) was needled as the reinforcement of the composites. The obtained carbon fiber fabric has a density of 0.52 g/cm3 and fiber volume fraction of about 32 vol%, which was first deposited with (PyCSiC)3 interphase by CVI process. And then, a Cf/ZrC-C nano-porous preform was prepared by impregnation of ZrC-C sol, followed by heat treatment. Finally, the Cf/ZrC-SiC composite was obtained by RMI of Si melt into the nano-porous Cf/ZrC-C preform at 1500 ℃ in the vacuum. Details of the fabrication process can be found in our previous work [20]. The PIP-Cf/ZrC-SiC composite was fabricated by conventional precursor infiltration and pyrolysis method for comparison.
RA =
l 0 − lt t
(2)
where RA is linear recession rate (μm/s), l0 and lt are thickness (μm) of composites before and after ablation. Specific heat and thermal diffusivity were measured by laser flash thermal conductivity meter (NETZSCH LFA467 HT, Germany). The size of the sample used for thermal diffusivity measurement is Φ10 mm × 2 mm. Thermal conductivity was calculated according to the equation below:
λ = αCp ρ
2.2. Test and characterization
(3) −1
−1
where λ and α are thermal conductivity (W⋅ m ⋅ K ) and thermal diffusivity (mm2⋅s−1), Cp and ρ are specific heat (J⋅ g−1⋅ K−1) and density (g⋅ cm-3) of the composites. The densities and open porosities of the composites were measured by the Archimedes’ method using deionized water as immersing medium. Phase compositions of the composites before and after
The as-fabricated composites were processed into a rectangular shape (˜60 mm × 30 mm). The plasma torch (Plasmer Jet, A-2000, Sulzer Metco, Switzerland) was used to test the ablation resistance of the specimens. The heat fluxes were 3.01 and 4.02 MW/m2, and the testing time was 150 s and 60 s respectively. The temperature of the 4618
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ablation test were determined by X-ray diffraction (XRD, Ultima IV, Rigaku Corporation, Kyoto, Japan) using Cu-Kα radiation. Microstructural features of the samples were characterized by scanning electron microscopy (SEM, S4800, Tokyo, Japan) combined with energy dispersive spectroscopy (EDS, Oxford Instruments, UK). Based on the different grayscale of ZrC and SiC in SEM image, the volume fractions of ZrC and SiC in composites were determined by ImageJ® freeware (https://imagej.nih.gov/ij/) through image analysis. To reduce the error of image analysis, at least 10 images were taken in random regions of the samples and used for statistical analysis. 3. Result and discussion 3.1. Microstructure and compositions The as-fabricated RMI-Cf/ZrC-SiC composite has a density of 2.64 g/ cm3 and a porosity of 3.49%, whereas the density and open porosity of the PIP-Cf/ZrC-SiC composite are 2.18 g/cm3 and 13.53% respectively. Fig. 1 illustrates the polished cross-sections of the composites. In agreement with the density measurement, RMI-Cf/ZrC-SiC shows a highly dense microstructure, with almost no obvious pores. Three different phases present in the matrix of the composites: bulk ZrC, white needle-shaped ZrSi2 and continuous dark grey SiC, as indicated in Fig. 1c according to EDS analysis. As reported before, the needle-shaped ZrSi2 is formed via the reaction between ZrC and Si during the cooling stage of RMI [20]. It can also be observed that fibers and interphase are well protected in this RMI-Cf/ZrC-SiC composite, without obvious fibers/ interphase erosion (Fig. 1e). In comparison, the PIP-Cf/ZrC-SiC composite presents a completely different microstructure. Quite a few small pores can be observed on the polished cross-section of the PIP-Cf/ ZrC-SiC composite, where the ZrC and SiC show a uniform distribution with fine particle size (Fig.1d and g). The volume fractions of ZrC and SiC in the two composites are determined based on image analysis and illustrated in Table 1. SiC is the main phase component in both composites. But the SiC content in the PIP-Cf/ZrC-SiC is a little bit less than that in the RMI-Cf/ZrC-SiC composite. Fig. 2 displays the XRD patterns of the as-fabricated composites, where the results are consistent with the SEM and EDS analysis. The RMI-Cf/ZrC-SiC composite is mainly composed of ZrC, SiC and ZrSi2, while the PIP-Cf/ZrC-SiC composite is composed of ZrC and SiC phases. In addition, the wide diffraction peak, located at around 26°, corresponds to carbon fibers of the composites.
Fig. 2. XRD patterns of composites before ablation.
the measured density, thermal diffusivity and specific heat, the calculated thermal conductivity of RMI-Cf/ZrC-SiC composite decreases gradually from 21.26 W⋅ m-1⋅ K-1 at room temperature to 15.43 W⋅ m1 ⋅ K-1 at 1200 °C. The thermal conductivity of RMI-Cf/ZrC-SiC composite is quadruple higher than that of the PIP-Cf/ZrC-SiC composite (Fig. 3c). The high thermal conductivity of the RMI-Cf/ZrC-SiC composite can be attributed to its low porosity (3.49%) and continuous phase distribution. The relation between thermal conductivity and porosity can be expressed roughly by the following equation [26].
λ = λs (1 − p)
Thermal diffusivity, specific heat and thermal conductivity of the composites were measured in the temperature range of RT-1200 °C. The results are shown in Fig. 3. The thermal diffusion coefficient is a typical physical quantity that describes how fast the heat is transferred in materials. The thermal diffusion coefficient of the as-fabricated RMI-Cf/ ZrC-SiC composite is 12.06 mm2·s−1 at room temperature, which decreases to 4.20 mm2·s−1 at 1200 °C. Specific heat is used to evaluate the energy storage capacity of materials. As shown in Fig. 3b, the specific heat of the RMI-Cf/ZrC-SiC composite increases gradually with the increase in temperature and reaches 1.47 J⋅ g-1⋅ K-1 at 1200 °C. Based on
3.3. Ablation-resistant property and mechanisms Fig. 4 illustrates the macroscopic features of the composites after ablation under different heat fluxes. The ablated surface can be divided into central ablation area and transition area. This typical ablation morphology is caused by the synergic effects of heat, oxidation and mechanical scouring. There’s no distinct ablation damage on the RMICf/ZrC-SiC composite (Fig. 4a-b). As shown in Fig. 4a, only a small amount of oxide spatters out from the ablation center after ablation by the high heat flux (4.02 MW/m2). In comparison, the PIP-Cf/ZrC-SiC composite is damaged remarkably and a specific ablation center can be clearly observed (Fig. 4c-d).
Table 1 Parameters of the RMI-Cf/ZrC-SiC and the PIP- Cf/ZrC-SiC composites.
RMI-Cf/ZrCSiC PIP- Cf/ZrCSiC
Density (g/ cm3)
Open porosity (%)
Phase volume fraction (%) ZrC
ZrSi2
SiC
Cf
2.64
3.49
6.69
6.76
50.52
32.54
2.18
13.53
13.22
N/A
41.20
32.05
−1
where λ is the thermal conductivity (W⋅ m ⋅ K ), λs is the thermal conductivity (W⋅ m−1⋅ K−1) of fully dense material, and p is the porosity (%) of the material. Assuming the RMI-Cf/ZrC-SiC composite has the same porosity (13.53%) as the PIP-Cf/ZrC-SiC composite, the thermal conductivity with 13.53% of porosity in the RMI-Cf/ZrC-SiC composite is calculated to be 19.05 W⋅ m−1⋅ K−1 at room temperature and 13.83 W⋅ m−1⋅ K−1 at 1200 ℃. It can be found that the estimated thermal conductivity values are also higher than those of the PIP-Cf/ZrC-SiC composite. This indicates that in addition to its low porosity, other factors also contribute to the high thermal conductivity of the RMI-Cf/ZrC-SiC composite. Generally, heat is conducted mainly through phonon motion of the solid framework (crystal lattice vibration) in composites [27]. Because of the dispersed distribution characteristics of ZrC particles in the PIPCf/ZrC-SiC composite (Fig. 1g), the SiC phase is separated into small pieces, which remarkably increases the number of phase boundaries. Consequently, the abundant phonons scatter at phase boundaries and high porosity (13.53%) leads to the low thermal conductivity of the PIP-Cf/ZrC-SiC composite.
3.2. Thermophysical properties
Samples
(4) −1
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Fig. 3. Thermal diffusivity (a), specific heat (b) and thermal conductivity (c) of RMI-Cf/ZrC-SiC and PIP-Cf/ZrC-SiC composites.
ablation resistance of the composites. As described in Fig. 6a, reactions (Eq. (7) and (8)) are more thermodynamically favorable at a lower temperature. Therefore, a continuous SiO2-ZrO2 layer tends to form on the ablated surface at the starting stage of the ablation test. With the increase of temperature, the chemical process of Eq. (5) and (6) becomes more active, and a mass of gaseous products (SiO and CO) are generated. The evaporation of SiO and CO results in the formation of a lot of pores in the SiO2-ZrO2 layer. The difference in microstructure of these two composites also gives influence to their oxidation behavior. The bulk ZrC (Fig. 1f) in the RMICf/ZrC-SiC composite is oxidized into continuous ZrO2 which dissolves in glassy SiO2. The continuous ZrO2 acts as a net structure to prevent SiO2 collapse from flame scouring and reduces the evaporation of SiO2 (Fig. 7a). In comparison, the dispersed distribution of ZrC (Fig. 1g) in PIP-Cf/ZrC-SiC is oxidized into dispersed ZrO2 which may has less pinning effect for glassy SiO2. So, the glassy SiO2 is easily scoured by flame (Fig. 7b). Fig. 6b shows the XRD patterns of the composites at ablation center. m-ZrO2 and SiC are detected in both the RMI-Cf/ZrC-SiC and the PIP-Cf/ ZrC-SiC composites at different heat flux conditions. It can be found that the diffraction peaks intensity of ZrO2 is relatively lower for the composites ablated at a heat flux of 3.01 MW/m2 due to limited ZrO2 formed under this condition. However, SiO2 is not detected in Fig. 6b. After ablation, the surface temperature drops sharply and the SiO2based glassy phase commonly stays in an amorphous state under this condition. As a result, no diffraction or diffusion peaks of the glassy phase are detected in the XRD spectra. This phenomenon has also been well reported in literature [4,28–30]. The existence of SiC indicates that SiC matrix is not completely covered by the glassy oxide layer. Some cracks and pores formed on the oxide layer during ablation (Fig. 7a and b), leading to the exposure of SiC. As shown in Fig. 7c and d, it can be observed that most of the sample surface is well covered by glassy layer at the ablated center, particularly for the RMI- Cf/ZrC-SiC. In order to understand the ablation behavior carefully, further observation was performed on the ablation area of the composites ablated at a heat flux of 4.02 MW/m2 for 60 s. The results are shown in Fig. 7.
The calculated ablation recession rates are plotted in Fig. 5. The mass and linear ablation rates of both composites at 4.02 MW/m2 are much higher than those of the composites at 3.01 MW/m2, indicating the composites are easier ablated at high heat flux, even tested for a shorter period (Fig. 5). The mass and linear ablation rates of the RMICf/ZrC-SiC composite decline by 39.02% and 98.07% respectively at 4.02 MW/m2 compared with the PIP-Cf/ZrC-SiC composite. The mass and linear ablation rates of the RMI-Cf/ZrC-SiC composite at 3.01 MW/ m2 are 33.53% and 98.71% smaller than those of the PIP-Cf/ZrC-SiC composite (Fig. 5). Therefore, RMI-Cf/ZrC-SiC composite shows a much better ablation-resistant property at two tested heat flux conditions. The flame temperature at high heat flux is about 3000 °C, which is much higher than that of low heat flux (˜2300 °C). Thus, the composites ablated at 4.02 MW/m2 have a higher surface temperature compare with the case of 3.01 MW/m2 (Table 2). Moreover, due to the much higher thermal conductivity, the heat on the ablated surface of the RMICf/ZrC-SiC composite can be transmitted and dissipated timely during the ablation, which reduces the surface temperature. As shown in Table 2, the surface temperature of the RMI-Cf/ZrC-SiC composite is 130–210 °C lower than that of the PIP-Cf/ZrC-SiC composite at the same ablation condition, which weakens the thermal damage of the RMI-Cf/ ZrC-SiC composite. Generally, the main reactions of the Cf/ZrC-SiC composites during the high-temperature ablation include: C + 0.5O2 (g) = CO (g)
(5)
SiC + O2 (g) = SiO (g) + CO (g)
(6)
SiC + 1.5O2 (g) = SiO2 + CO (g)
(7)
ZrC + 1.5O2 (g) = ZrO2 + CO (g)
(8)
The changes of standard Gibbs free energy of the above reactions are calculated as a function of temperature and plotted in Fig. 6a. The formation of a continuous SiO2-ZrO2 layer on the ablated surface can protect the composites from further oxidation and ablation. However, the evaporation of gaseous products during ablation is harmful to the 4620
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Fig. 4. Macroscopic features of the composites after ablation at different heat fluxes: (a) RMI-Cf/ZrC-SiC ablated at 4.02 MW/m2 for 60 s; (b) RMI-Cf/ZrC-SiC ablated at 3.01 MW/m2 for 150 s; (c) PIP-Cf/ZrC-SiC ablated at 4.02 MW/m2 for 60 s; (d) PIP-Cf/ZrC-SiC ablated at 3.01 MW/m2 for 150 s.
composite. By contrast, the dispersed distribution of ZrC in PIP-Cf/ZrCSiC composite divides SiC phase into small pieces along with a mass of small pores (Fig. 1d and g), which provide oxygen diffusing paths and result in the fragmentized oxide layer (Fig.7d). Fig. 7 e–f show the microstructure of the transition area on the ablated surface. As the temperature of the transition area is much lower than that of the ablation center, both samples are covered by the continuous glassy oxide layer, without evident signs left by massive gaseous product. As discussed above, the ablation behavior can be affected by microstructure, thermophysical properties and the formed oxide layer, etc. of the composites. Considering their interactions, the different ablationresistant property and mechanism between the RMI-Cf/ZrC-SiC and the PIP-Cf/ZrC-SiC composites can be summarized as Fig. 8:
After ablation test, the ablation center of the RMI-Cf/ZrC-SiC composite is covered by a continuous oxide layer, as shown in Fig. 7a and c. The EDS analysis (Fig. 7g) reveals that the continuous oxide layer is composed of SiO2 glass with ZrO2 particles embedding in it. Because SiO2 playing as a bonding agent, oxygen cannot easily diffuse through the continuous SiO2-ZrO2 layer. The continuous SiO2-ZrO2 layer plays a significant role in improving the ablation-resistant property as it can separate the surface of the composites from the ablation atmosphere and protect the underlying composites from damage. In comparison, the surface temperature of the PIP-Cf/ZrC-SiC composite is as high as 2100 °C during ablation test at a heat flux of 4.02 MW/m2, which is 210 °C higher than that of the RMI-Cf/ZrC-SiC composite (Table 2). As a result, SiO and CO are generated heavily by reactions (Eq. (5) and (6)) and escaped quickly through the surficial oxide layer. Meanwhile, SiO2 cannot survive due to its serious evaporation at such high temperature. Consequently, the ablation center of the PIP-Cf/ZrC-SiC composite exhibits a loose powdery microstructure, as shown in Fig. 7b and d. As a result, the protective effect of the surficial oxide layer on the PIP-Cf/ ZrC-SiC composite is weakened at such condition. On the other side, under the surface oxide layer, a porous layer presents in both the RMICf/ZrC-SiC and the PIP-Cf/ZrC-SiC composites, where the pores form due to the recession of SiC and carbon fibers during ablation. The difference in ablation-resistant property between the two composites is also related to their microstructure and phase constituent. The dense and continuous microstructure of the RMI-Cf/ZrC-SiC composite can effectively retard oxygen from diffusing into the interior of the composite and further slow down the ablation recession of the
(i) At the beginning of ablation, the exterior ZrC and SiC are oxidized and a continuous SiO2-ZrO2 glassy layer forms and covers on the sample surface. (ii) With the ablation process going on and the surface temperature increasing, the production and recession of SiO2-ZrO2 layer reach a dynamic equilibrium state. Benefiting from the low porosity and high thermal conductivity of RMI-Cf/ZrC-SiC composite, the heat is timely delivered from the ablation surface. Ablation center temperature is kept below 1900 °C even though the flame temperature is as high as 3000 °C at a heat flux of 4.02 MW/m2. Consequently, the continuous SiO2-ZrO2 protective layer is retained on the composite surface during the ablation process, leading to an excellent
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Fig. 5. Linear recession rate and mass recession rate under different ablation condition.
structure. (iv) The loose powdery structure of the surficial oxide cannot provide effective protection for the PIP-Cf/ZrC-SiC composite. Consequently, the recession of inside matrix and carbon fibers is aggravated, and the damage is going inward.
Table 2 Parameters of the ablation test. Ablation condition Heat flux (MW/m2)
Time (s)
4.02
60 s
3.01
150 s
Samples
Surface temperature (℃)
RMI-Cf/ZrC-SiC PIP- Cf/ZrC-SiC RMI-Cf/ZrC-SiC PIP- Cf/ZrC-SiC
˜1890 ˜2100 ˜1700 ˜1830
4. Conclusion Cf/ZrC-SiC-based composite is fabricated by an improved reactive melt infiltration (RMI). The microstructure and ablation-resistant property of the composite are investigated comprehensively, and compared with conventional PIP-Cf/ZrC-SiC composite. The RMI-Cf/ ZrC-SiC composite has low porosity (3.49%) and high thermal conductivity (21.3 W⋅ m−1⋅ K−1 at room temperature, 15.4 W⋅ m−1⋅ K−1 at 1200 °C), which is beneficial for transferring the heat timely and reducing the surface temperature of the composite during ablation. As a
ablation-resistant property for the RMI-Cf/ZrC-SiC composite. (iii) For PIP-Cf/ZrC-SiC composite, its limited thermal conductivity causes a sharp increase of the surface temperature. SiO2 evaporated quickly along with a heavy generation of SiO and CO gas. As a result, ZrO2 is left on the surface and forms a loose powdery
Fig. 6. (a) Changes of standard Gibbs free energy of reaction (5)-(8) as a function of temperatures calculated by HSC 6.0 software [31,32]; (b) XRD patterns of the composites at ablation center. 4622
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Fig. 7. SEM images of ablated surfaces and cross-sections at a heat flux of 4.02 MW/m2 for 60 s. (a, c) ablation center of RMI-Cf/ZrC-SiC; (b, d) ablation center of PIPCf/ZrC-SiC; (e, f) transition area of ablated RMI-Cf/ZrC-SiC and PIP-Cf/ZrC-SiC; (g) EDS analysis of the spots marked in (a), (b).
ZrC-SiC composite. During ablation, a continuous glassy SiO2-ZrO2 layer can be formed and cover on the ablation surface to hinder the ingression of oxygen into the RMI-Cf/ZrC-SiC composite. The layer also isolates the underlying materials from the plasma flame.
result, the RMI-Cf/ZrC-SiC composite presents an excellent ablationresistant property at two tested heat fluxes. The linear and mass recession rates of the RMI-Cf/ZrC-SiC composite decline by 98.71% and 33.53% at a heat flux of 3.01 MW/m2, and they decrease by 98.07% and 39.02% at a heat flux of 4.02 MW/m2 compared with the PIP-Cf/
Fig. 8. Schematic diagram of the ablation mechanism for Cf/ZrC-SiC composites. 4623
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