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SiC composites under molten fluoride salt environment
Journal of Nuclear Materials 487 (2017) 43e49
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Journal of Nuclear Materials journal homepage: www.elsevier.com/locate/jnucmat
The corrosion behavior of CVI SiC matrix in SiCf/SiC composites under molten fluoride salt environment Hongda Wang a, b, c, Qian Feng d, Zhen Wang a, b, *, Haijun Zhou a, b, Yanmei Kan a, b, Jianbao Hu a, b, Shaoming Dong a, b, ** a
Structural Ceramics and Composites Engineering Research Center, Shanghai Institute of Ceramics, Chinese Academy of Sciences, Shanghai 200050, China State Key Laboratory of High Performance Ceramics and Superfine Microstructure, Shanghai Institute of Ceramics, Chinese Academy of Sciences, Shanghai 200050, China c School of Graduate, University of Chinese Academy of Sciences, Beijing 100049, China d Analysis and Testing Center, Donghua University, Shanghai 201600, China b
a r t i c l e i n f o
a b s t r a c t
Article history: Received 24 July 2016 Accepted 2 February 2017 Available online 5 February 2017
High temperature corrosion behavior and microstructural evolution of designed chemical-vaporinfiltrated SiC matrix in SiC fiber reinforced SiC ceramic matrix composites in 46.5LiF-11.5NaF-42.0KF (mol. %) eutectic salt at 800 C for various corrosion time was studied. Worse damage was observed as extending the exposure time, with the mass loss ratio increasing from 0.716 wt. % for 50 h to 5.914 wt. % for 500 h. The mass loss rate showed a trend of first decrease and then increase with the extended corrosion exposure. Compared with the near-stoichiometric SiC matrix layers, the O-contained boundaries between deposited matrix layers and the designed Si-rich SiC matrix layers were much less corrosion resistant and preferentially corroded. Liner relationship between the mass loss ratio and the corrosion time obtained from 50 h to 300 h indicated that the corrosion action was reaction-control process. Further corrosion would lead to matrix layer exfoliation and higher mass loss ratio. © 2017 Published by Elsevier B.V.
1. Introduction Silicon carbide (SiC) continuous fiber-reinforced SiC matrix (SiCf/SiC) composites are promising, commercially available materials for nuclear reactors applications. The advanced properties, including the excellent high temperature mechanical performance [1,2] and low activation properties [3e5] inherent to SiC, make the SiCf/SiC composites to be an attractive candidate material for structural and functional components in nuclear reactors. Nowadays, significant progress has been made towards the goal of understanding the stability of SiCf/SiC composites under irradiation condition, including the fundamental properties such as strength and dimensional stability at high temperature [6,7]. However, the compatibility of SiC-based materials with coolant has been
* Corresponding author. Shanghai Institute of Ceramics, Chinese Academy of Sciences, No. 1295, Dingxi Road, Changning District, Shanghai 200050, China. ** Corresponding author. Shanghai Institute of Ceramics, Chinese Academy of Sciences, No. 1295, Dingxi Road, Changning District, Shanghai 200050, China. E-mail addresses: [email protected] (Z. Wang), [email protected] (S. Dong). http://dx.doi.org/10.1016/j.jnucmat.2017.02.007 0022-3115/© 2017 Published by Elsevier B.V.
demonstrated seldomly up to now, especially the database of corrosion behavior of SiCf/SiC composites in variety of coolants at high temperature. More recently, researchers [8e10] focus on the excellent compatibility of SiCf/SiC composites with molten fluoride salt (e.g. 46.5 mol % LiF- 11.5 mol % NaF-42.0 mol % KF eutectic molten salt, i.e. FLiNaK salt), one kind of candidate coolants used in the new Generation IV Nuclear Reactor [11,12]. This unique character make SiCf/SiC composites be one of the alternative materials for heat-exchanger, control rod and the other structure components in the new nuclear reactor concepts, such as the Molten Salt Reactor (MSR) and the Very High Temperature Reactor (VHTR) [13e15]. The corrosion of material in the liquid fluoride salt at high temperature is related to its composition. Free Si in C/SiSiC composite ceramic materials prepared by melt infiltration (MI) processing would be corroded selectively when exposed to the FLiNaK salt at 850 C for 500 h, while the SiC phase and remained pure C phase are relatively less attacked. This is qualitatively understood with the respect of the Gibbs free energy of fluoride formation per F2 molecule of the Si and SiC phases [16]. SiC fiber with higher oxygen content suffers worse corrosion in molten FLiNaK salt at
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700 C for 15 days than SiC fiber contained much fewer oxygen, indicating that the oxygen in SiC accelerates the corrosion of SiC in fluoride salt [10]. To meet the stricter service requirements in the highly corrosive liquid fluoride salt, the nuclear grade SiCf/SiC composites should be composed of high-purity, near-stoichiometric, and dense fibers and matrix. Various processes have been employed for SiCf/SiC composites fabrication, including chemical vapor infiltration (CVI), polymer impregnation and pyrolysis (PIP), reaction sintering (RS) and NanoInfiltration and Transient Eutectic phase (NITE) processing [17e20]. Among the various fabrication processes, the chemical vapor infiltration (CVI) processing is attractive for the fabrication of SiCf/ SiC composites because it yields highly crystalline, nearstoichiometric SiC matrix and the fiber damage during processing is minimized, which are widely employed to prepare the nuclear grade SiCf/SiC composites. However, the uniformity of deposited SiC matrix composition during CVI processing is determined by the temperature evenness in the furance, especially in the fabrication of the large-scale component. Owing to the existence of temperature gradient, the matrix located in the higher temperature region would be high-purity, near-stoichiometric SiC, and the matrix obtained in the lower temperature region would be composed of codeposited free Si and SiC [21]. The co-deposited free Si would affect the corrosion resistance performance of SiCf/SiC composites in liquid fluoride salt. Nevertheless, the effect of non-uniformity of matrix composition obtained in CVI processing to the corrosion behavior of SiCf/SiC composites has been report rarely so far, which is overwhelmingly important for the safe application of the material. In this work, SiCf/SiC composites with designed matrix, which is composed of Si-rich SiC layers and near-stoichiometric SiC layers, are prepared by controlling the CVI temperature to simulate the non-uniform SiC matrix caused by temperature gradient. The corrosion behavior of these designed SiC matrix in molten fluoride salt are studied. The microstructural evolution during the corrosion process, the corrosion kinetics at 800 C and the underlying corrosion processing are experimentally and theoretically investigated.
scale component. The fiber volume fraction and porosity of asreceived composites were ~40% and ~15%. The composites were cut into 35(long) 4(wide) 2(thick) mm3 bar for the corrosion experiment. Each corrosion experiment group contains 5 bars. The experimental setup of the corrosion test is schematically shown in Fig. 1. Due to the good chemical compatibility of graphite with fluoride molten salt, the corrosion test was carried out in a graphite crucible. Before corrosion, the graphite crucible and SiCf/ SiC samples were heat-treated in furnace at 700 C for 10 h to eliminate possible adsorptions. Therein, SiCf/SiC samples were fixed on a graphite support and dipped into ~80 g 46.5LiF - 11.5NaF - 42.0KF (mol. %, FLiNaK) fluoride eutectic molten salt (Shanghai Institute of Organic Chemistry, Chinese Academy of Sciences, Shanghai), which contained ~120 ppm O content. Subsequently, the crucibles were furnace heated at 800 C for 50 h, 100 h, 200 h, 300 h or 500 h, respectively. Here, it should be mentioned that all the above operations were conducted in a glove box under the protection of high purity argon gas. The atmosphere in the glove box was controlled at O2 <0.1 ppm and moisture <0.1 ppm during the corrosion process. After the high temperature corrosion process in fluoride eutectic molten salt, the SiCf/SiC composites samples were picked out from the crucible and immersed into 1 mol/L Al(NO3)3 solution to remove the fluoride salt solid adhering on the exterior and interior surfaces of the sample [23]. Finally, the samples were cleaned by ultrasonicating in deionized water, then dried in a vacuum drying chamber. The samples were weighed by precision balance (BT-25-S, 0.01 mg, Sartorius) before and after the corrosion test, the mass loss ratio, and mass loss rate, were applied to characterize the corrosion resistance property of the composites and corrosion kinetics. Phase compositions of the samples were analyzed by X-ray diffraction (XRD, D8 Advance, Bruker AXS Co. Ltd.) at room temperature, using CuKa radiation (l ¼ 0.154 nm) at a scanning 2q angles range of 10 e80 . The microstructure and chemical composition of samples before and after the corrosion test were characterized by scanning electron microscope (SEM, Quanta-250, FEI) and energy dispersive spectrometer (EDS, Oxford), respectively. 3. Results
2. Experimental The materials evaluated are two-dimensional woven fabrics of SiC fibers reinforced chemical vapor infiltrated SiC matrix composites. The used SiC fibers, known as KD-II SiC fibers, were provided by the National University of Defense Technology, Changsha, Hunan Province, China. The basic properties are shown in Table 1. More key characteristics of these SiC fibers, including the corrosion resistance of fibers in FLiNaK molten salt, have been reported elsewhere [10,22]. In advance of the CVI process, a pyrolytic carbon (PyC) monolayer about 100 nm in thickness was deposited on the SiC fiber surface through chemical vapor deposition (CVD). Using the Methyltrichlorosilane (MTS) as the raw material and H2 as the carrier gas, the SiCf/SiC composites with designed Si-rich SiC layers and near-stoichiometric SiC layers in matrix were prepared by controlling the deposited temperature via CVI processing. The purpose of preparation Si-rich SiC layers in SiC matrix was to simulate the non-uniformity of SiC matrix composition caused by temperature gradient in the CVI process for preparation the large
Fig. 2 shows mass loss ratio and mass loss rate of the samples after corrosion at 800 C for vary times in the molten FLiNaK salt. The mass loss ratio increased from 0.716 wt. % to 5.914 wt. % as the corrosion time was prolonged from 50 to 500 h, indicating a cumulative corrosion damage of the sample. The fitting line showed that the samples' mass loss ratio, ml, is as a function of the corrosion time, t, at 800 C. Linear relationship ml ¼ f(t) was observed when t was among 0e300 h. However, the mass loss ratio of samples after corrosion for 500 h deviated from the fitting line seriously. Furthermore, with the extended corrosion exposure, the mass loss rate showed a trend of first decrease and then increase and could be divided into three stages: the decelerated (I), constant rate (II) and accelerated (III) stages. The surface morphologies of the SiCf/SiC samples before and after corrosion are shown in Fig. 3. Compared with the flat surface of samples before corrosion, the sample surface became obviously rougher after corrosion, and matrix damage could be distinctively observed. For the sample corroded for 100 h, parts of the matrix
Table 1 General properties of the as-received KD-II SiC fibers [22]. Properties
C/Si mole ratio
Oxygen (wt %)
Density (g/cm3)
Diameter (mm)
Strength (GPa)
Modulus (GPa)
KD-II
1.41
1.3
2.7
12.1
3.0
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H. Wang et al. / Journal of Nuclear Materials 487 (2017) 43e49
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Fig. 1. Schematic illustration of (a) the experimental assembly and (b) the graphite crucible setup used for molten salt corrosion test.
Fig. 2. Mass loss ratio and mass loss rate of samples corroded at 800 C for various time. The minus sign of mass loss rate represents mass decrement.
layer on its surface were exfoliated as indicated by the arrows in Fig. 3(b). When the corrosion time was prolonged to 500 h, more intensive matrix corrosion and exfoliation occurred, which was in good agreement with the large increase in the mass loss ratio of the composites. At this time, SiC fibers with remnant thin SiC deposition layer exposed on the sample surface, and layer-exfoliation of the SiC layers were frequently observed (indicated by black arrow in Fig. 3). The cross section images of the samples before and after corrosion are shown in Fig. 4. Matrix with multilayer structure was observed in samples before and after corrosion. For the sample before corrosion, two types of designed matrix layers were found in the composites, the light gray matrix layer and deep gray matrix
layer. These matrix layers were separated with interlayer boundaries (as shown in white box in Fig. 4(a)). EDS analysis results shown in Table 2 revealed that the deep gray matrix layer showed atomic ratio of Si/C z 1.0 (point 1 in Fig. 4(a)), while the light gray one was significantly Si-rich (point 2 in Fig. 4(a)). This indicated that the deep gray matrix layer should be mainly composed of SiC, whereas a large amount of free Si should exist in the designed light gray matrix layer. Linescan analysis of interlayer boundary area in deposited matrix shown in the white box in Fig. 4(a) suggested stronger intensity of oxygen element in the boundary than matrix. In another word, the interlayer boundaries contained more oxygen than deposited matrix. After being exposed the fluoride molten salt at 800 C for 100 h (Fig. 4(b)), the light gray matrix layers and interlayer boundaries nearly disappeared, due to the molten salt corrosion. This suggested the poorer corrosion resistance of the Si-rich SiC layers and the interlayer boundaries in comparison with the SiC layers. EDS analysis of the corrosion regions (point 3 and 4 in Fig. 4(c)) revealed significant silicon losses, which made the chemical composition of these regions shift from silicon rich to carbon rich. Further corrosion of the composites led to the formation of vesicular structured layers in the matrix as the corrosion time was prolonged to 500 h, which contributed to the significant increment of the mass loss ratio. On another aspect, the oxygen content in corrosion area got higher than that in deposited layers of virgin samples. Fig. 5 shows the XRD patterns of the SiCf/SiC composites before and after corrosion at 800 C for 100 and 500 h. The samples before and after corrosion for 100 h showed quite similar XRD patterns, where diffraction peaks solely assigned to b-SiC were detected. While for the sample after corrosion for 500 h, new diffraction peak corresponding to Graphite-2H was identified in its XRD patter in addition to the peaks of b-SiC. Considering the high C content at the corrosion area obtained by EDS analysis, Graphite-2H may be the
Fig. 3. SEM images of the surface morphology of samples (a) before corrosion, and after corrosion at 800 C for (b) 100 h and (c) 500 h.
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Fig. 4. SEM images of the cross sections of the samples: (a) before corrosion, (b) corroded for 100 h, (c) higher magnification of area in white box in (b), and (d) corroded for 500 h.
Table 2 EDS Analysis of the local chemical composition on cross sections of the SiCf/SiC ceramic matrix composites (at. %). Point
1
2
3a
4a
C Si O
53.02 45.42 1.56
29.44 70.56 e
60.85 30.48 5.85
61.50 28.63 8.00
a
The F, K etc., elements of remained fluoride molten salt, with low content were not listed in the table.
corrosion product of b-SiC in the molten salt. 4. Discussion 4.1. Corrosion kinetics Solid reaction consists of many simple physical and chemical processes, such as chemical reaction, diffusion, crystallization, and so on [24]. As shown in Fig. 6(a), the solid reaction forms a reaction
Fig. 5. XRD patterns of the SiCf/SiC composites (a) before and (b) after corrosion at 800 C for 100 and 500 h.
H. Wang et al. / Journal of Nuclear Materials 487 (2017) 43e49
47
Fig. 6. Schematic drawing of solid corrosion models.
layer on the contact surface, for example, SiO2 layer formed by passive oxidation of Si or SiC [25e27]. The kinetics of solid reaction would be controlled by the slowest process, especially the reaction process and diffusion process [28]. For the corrosion of Si and SiC in liquid fluoride salt, the kinetics could be expressed by the equation below [24]:
n¼
1 kcn0
1 þ Dcx 0
(1)
where n is corrosion rate, c0 is the concentration of corrosion media, k is the corrosion reaction constant, n is the order of reaction, D is the diffusion constant of corrosion agent, and x is the thickness of formed reaction layer. When SiC or Si is dipped into high temperature molten fluoride salt, SiF4 is believed to be the corrosion production, which would dissolve into the liquid fluoride salt [29], and no dense layer would be obtained on the surface of SiCf/SiC composites in liquid fluoride salt (shown in Fig. 6(b)). As a result, there is no diffusion process during corrosion (x ¼ 0, D ¼ 0). In another word, the corrosion kinetics is controlled by reaction process
n ¼ kcn0
(2)
Moreover, the graphite, one of the corrosion production showed in XRD patterns, would have no contribution to the mass loss, which suggests that the source of mass loss would be the loss of Si from both free Si and SiC. The relationship between mass loss, Dm, and corrosion time, t, could be expressed as follows:
Dm ¼ MSi $
Z
ndt ¼ MSi $
Z kcn0 dt ¼ MSi kcn0 t
(3)
with Dm being mass loss of samples, t the corrosion time, and MSi the molar mass of Si. Then the relationship between mass loss ratio, ml, and corrosion time, t, could be obtained:
ml ¼
Dm m
¼
MSi k$cn0 $t m
(4)
where ml and m are the mass loss ratio and the original mass of samples, respectively. In consideration of the corrosion in pure liquid fluoride salt, the concentration of corrosion media, c0, would be invariant. That is to say, the corrosion kinetics of SiCf/SiC composites in liquid fluoride salt concides with linear law. This is in consist with the liner relationship between mass loss ratio, ml, and the corrosion time, t, when t is ranged from 0 to 300 h. However, the mass loss rate, which corresponds to the corrosion rate that should be constant, shows the trend of decrease (Stage I in Fig. 2) first, and then being stable from 100 h to 300 h in Stage II. The higher mass loss rate at first is associated with the quick dissolution of oxidation layer on surface of SiCf/SiC composites and the high O-content boundaries between deposited matrix layers, which leads to larger mass loss in short time. As the depletion of the O-contained compositions, the corrosion reaction was shifted to the direct corrosion of matrix with
much lower rate. As a result, the mass loss rate would be dominated by the simple matrix corrosion with relatively stable corrosion rate (Stage II). On the other aspect, the mass loss ratio of the samples corroded for 500 h deviated from the fitting line seriously, and the mass loss rate increases slightly (Stage III). This would be associated to the worse damage after corrosion in FLiNaK salt for 500 h. As shown in Fig. 3(c), damaged fibers with no SiC deposition layer in the samples corroded for 500 h exposed because of the exfoliation of deposition matrix layers. This may be imputed to that liquid FLiNaK salt would erode the Si-rich SiC layer and boundaries between the deposition layers at first. After being destroyed the Si-rich layers and boundaries, the bonding force between deposition layers is gone, which leads to the more severe layers exfoliation and higher mass loss than that caused by corrosion reaction only. 4.2. Microstructural evolution The microstructural evolution of the designed CVI matrix before and after corrosion in liquid LiNaKF salt at 800 C shown in Fig. 4 clearly demonstrates that there are two kinds of corrosion mechanism during the erosion of designed matrix. These two kinds of mechanism are totally different. In view of the non-uniformity of deposited matrix caused by temperature gradient in the CVI furnace [30], especially in the large-scale CVI instrument for large-scale components preparation, Si-rich SiC layers are designed and introduced into the matrix besides near-stoichiometric SiC layers by controlling CVI parameters [31] to simulate the non-uniform CVI SiC matrix. The microstructure of SiCf/SiC composites before and after corrosion at 800 C shown in Fig. 4 indicated that these two kinds of designed matrix layers have different corrosion behaviors. The designed Si-rich SiC layers suffer worse corrosion than near-stoichiometric SiC layers. This is corresponding to the research of Jens Schmidt et al. [16] conducted at 850 C. As reported in this research, a qualitative characteristic of the corrosion potential of the phases made by the Gibb's free energy of fluoride formation per F2 molecule (DGf) of the Si and SiC phases; the more negative the DGf for a certain phase, the more this phase can be easily corroded. Calculation results show that the DGf of SiF4 formed from Si (727 kJ/mol) is more negative than from SiC (696 kJ/mol), suggesting that the Si has poor corrosion resistance than SiC at 850 C. This result implies that the co-deposited Si in matrix is the weak area and should be avoided to the maximum extent during CVI processing for preparing the nuclear grade SiCf/SiC composites. During the preparation of SiC matrix by CVI processing, repeated grinding of the surface as it seal off and reimpregnating are usually needed for decreasing the porosity and obtaining denser composites [32]. Before the next infiltration processing, the deposited layer surface obtained prior infiltration would be oxidized slowly. When subsequent infiltration is finished, the very thin oxidation layer is covered by another deposited SiC layer and forms the higheroxygen-content boundary between deposited SiC matrix layers. During corrosion, the higher oxygen content deteriorates the corrosion resistance performance of boundaries, which is identical
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H. Wang et al. / Journal of Nuclear Materials 487 (2017) 43e49
Fig. 7. Schematic diagrams showing the high temperature corrosion process of SiCf/SiC composites in fluoride molten salt.
to the results reported by Yang et al. [10]. Silicon oxide would be dissolved by fluoride, and the fluorine ion would attack the Si-O-Si bond and form Si-F bond [33], which leads to the selectively corrosion of layer boundaries. These corroded boundaries will be used as the penetration channels of liquid fluoride salt. As to the corrosion of SiC in fluoride molten salt, C seemed to be one of the corrosion products. The appearance of C maybe caused by the corrosion of SiC by either FLiNaK salt or the impurity, i.e. Fe and Ni, in fluoride salt, which needs to be studied further. 4.3. Corrosion mechanism Based on the above results, the corrosion process of the CVI SiC matrix of SiCf/SiC composites in high temperature fluoride molten salt can be schematically described as Fig. 7. First, the molten salt infiltrates into the sample through its surface defects, such as pores or cracks. The oxided surface and highe O-content interlayer boundaries are dissolved quickly. Therein, the Si-rich SiC matrix that may be introduced by temperature gradient will be preferentially corroded and will lead to the formation of fissure-like channels that allow the molten salt penetrating deeply into the matrix. After extended corrosion exposure, the Si-rich SiC layers are seriously corroded. Also, the bonding between deposited layers disappears due to the corrosion of boundaries between layers and the damage of Si-rich SiC layers. These corrosion damages result in the formation of pores in the matrix layers and the exfoliation of matrix layers around fibers. 5. Conclusions The high temperature corrosion behavior and microstructural evolution of SiCf/SiC composites with designed matrix in LiF-NaFKF eutectic molten salt at 800 C for vary corrosion time was studied. According to the experimental results, the following conclusions could be drawn: 1. The designed Si-rich SiC layers in matrix and O-contained boundaries between deposited layers are the weak area of the composites in terms of the corrosion resistance to fluoride molten salt. They are more liable to be corroded by the fluoride molten salt in comparison with the near-stoichiometric SiC layers. 2. The mass loss ratio induced by the molten salt corrosion shows an obvious increase with the extending of exposure time. The liner corrosion kinetics from 0 to 300 h is obtained during the corrosion of Si and SiC in FLiNaK at high temperature and is coincide with the reaction-control process. The deviation of 500 h samples is caused by the both reaction corrosion and SiC matrix layer exfoliation. 3. During the corrosion of the CVI SiC matrix in fluoride molten salt, the Si-rich SiC layers and O-contained boundaries of deposited layers are preferentially corroded, followed by the
corrosion of SiC layers. Then the molten fluoride salt penetrates into the matrix and leads to further damage. 4. Considering the much worse corrosion resistance performance of Si, parameters of CVI processing for fabrication of nuclear grade SiCf/SiC composites with better compatibility to fluoride salt should be controlled to avoid the co-deposition of free Si. Acknowledgement The authors gratefully acknowledge the financial support from the National Natural Science Foundation of China (No. 51502323) and the Natural Science Foundation of Shanghai, China (No. 14ZR1445800). References [1] M. Gomina, P. Fourvel, M.-H. Rouillon, J. Mater. Sci. 26 (7) (1991) 1891. [2] T.K. Jacobsen, P. Brøndsted, J. Am. Ceram. Soc. 84 (5) (2001) 1043. [3] K.Y. Donaldson, B.D. Trandell, Y. Lu, D. Hasselman, N. Maurer, J. Am. Ceram. Soc. 81 (6) (1998) 1583. [4] W. Kowbel, C. Bruce, K. Tsou, K. Patel, J. Withers, G. Youngblood, J. Nucl. Mater 283 (2000) 570. [5] N. Taylor, I. Cook, R. Forrest, C. Forty, W. Han, D. Ward, Fusion Eng. Des. 58 (2001) 991. [6] Y. Katoh, K. Ozawa, T. Hinoki, Y. Choi, L.L. Snead, A. Hasegawa, J. Nucl. Mater 417 (1) (2011) 416. [7] Y. Katoh, T. Nozawa, L.L. Snead, K. Ozawa, H. Tanigawa, J. Nucl. Mater 417 (1) (2011) 400. [8] X. He, J. Song, J. Tan, B. Zhang, H. Xia, Z. He, X. Zhou, M. Zhao, X. Liu, L. Xu, J. Nucl. Mater 448 (1) (2014) 1. [9] H. Nishimura, T. Terai, T. Yoneoka, S. Tanaka, A. Sagara, O. Motojima, J. Nucl. Mater 283 (2000) 1326. [10] X. Yang, M. Liu, Y. Gao, D. Zhang, S. Feng, H. Liu, G. Yu, G. Wu, M. Wang, X. Zhou, Corros. Sci. 103 (2016) 165. [11] D. Williams, L. Toth, K. Clarno, Assessment of Candidate Molten Salt Coolants for the Advanced High Temperature Reactor (AHTR), Oak Ridge National Laboratory, 2006. ORNL/TM-2006/12, 2006. [12] M. Seregin, A. Parshin, A.Y. Kuznetsov, L. Ponomarev, S. Mel’nikov, A. Mikhalichenko, A. Rzheutskii, R. Manuilov, Radiochemistry 53 (5) (2011) 491. [13] D.F. Wilson, Chemical Compatibility Issues Associated with Use of SiC/SiC in Advanced Reactor Concepts, Oak Ridge National Laboratory, 2012. ORNL/TM2012/546. [14] Y. Katoh, L.L. Snead, C.H. Henager Jr., A. Hasegawa, A. Kohyama, B. Riccardi, H. Hegeman, J. Nucl. Mater 659 (2007) 367e370. Part A. [15] Y. Katoh, D.F. Wilson, C.W. Forsberg, Assessment of Silicon Carbide Composites for Advanced Salt-cooled Reactors, Oak Ridge National Laboratory, 2007. ORNL/TM-2007/168 Revision. [16] J. Schmidt, M. Scheiffele, M. Crippa, P.F. Peterson, E. Urquiza, K. Sridharan, L.C. Olson, M.H. Anderson, T.R. Allen, Y. Chen, Int. J. Appl. Ceram. Technol. 8 (5) (2011) 1073. [17] T. Taguchi, T. Nozawa, N. Igawa, Y. Katoh, S. Jitsukawa, A. Kohyama, T. Hinoki, L. Snead, J. Nucl. Mater 329 (2004) 572. [18] J.-S. Park, A. Kohyama, T. Hinoki, K. Shimoda, Y.-H. Park, J. Nucl. Mater 367 (2007) 719. [19] Y. Katoh, M. Kotani, H. Kishimoto, W. Yang, A. Kohyama, J. Nucl. Mater 289 (1) (2001) 42. [20] S. Lee, J. Park, Y. Katoh, A. Kohyama, D. Kim, J. Lee, H. Yoon, J. Nucl. Mater 307 (2002) 1191. [21] A.I. Kingon, L.J. Lutz, P. Liaw, R.F. Davis, J. Am. Ceram. Soc. 66 (8) (1983) 558. [22] H. Araki, H. Suzuki, W. Yang, S. Sato, T. Noda, J. Nucl. Mater (Part 2) (1998) 258e263, 1540. [23] L. Olson, J. Ambrosek, K. Sridharan, M. Anderson, T. Allen, M. Corradini, Evaluation of material corrosion in molten fluoride salt, in: Presentation on
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The corrosion behavior of CVI SiC matrix in SiCf/SiC composites under molten fluoride salt environment Hongda Wang a, b, c, Qian Feng d, Zhen Wang a, b, *, Haijun Zhou a, b, Yanmei Kan a, b, Jianbao Hu a, b, Shaoming Dong a, b, ** a
Structural Ceramics and Composites Engineering Research Center, Shanghai Institute of Ceramics, Chinese Academy of Sciences, Shanghai 200050, China State Key Laboratory of High Performance Ceramics and Superfine Microstructure, Shanghai Institute of Ceramics, Chinese Academy of Sciences, Shanghai 200050, China c School of Graduate, University of Chinese Academy of Sciences, Beijing 100049, China d Analysis and Testing Center, Donghua University, Shanghai 201600, China b
a r t i c l e i n f o
a b s t r a c t
Article history: Received 24 July 2016 Accepted 2 February 2017 Available online 5 February 2017
High temperature corrosion behavior and microstructural evolution of designed chemical-vaporinfiltrated SiC matrix in SiC fiber reinforced SiC ceramic matrix composites in 46.5LiF-11.5NaF-42.0KF (mol. %) eutectic salt at 800 C for various corrosion time was studied. Worse damage was observed as extending the exposure time, with the mass loss ratio increasing from 0.716 wt. % for 50 h to 5.914 wt. % for 500 h. The mass loss rate showed a trend of first decrease and then increase with the extended corrosion exposure. Compared with the near-stoichiometric SiC matrix layers, the O-contained boundaries between deposited matrix layers and the designed Si-rich SiC matrix layers were much less corrosion resistant and preferentially corroded. Liner relationship between the mass loss ratio and the corrosion time obtained from 50 h to 300 h indicated that the corrosion action was reaction-control process. Further corrosion would lead to matrix layer exfoliation and higher mass loss ratio. © 2017 Published by Elsevier B.V.
1. Introduction Silicon carbide (SiC) continuous fiber-reinforced SiC matrix (SiCf/SiC) composites are promising, commercially available materials for nuclear reactors applications. The advanced properties, including the excellent high temperature mechanical performance [1,2] and low activation properties [3e5] inherent to SiC, make the SiCf/SiC composites to be an attractive candidate material for structural and functional components in nuclear reactors. Nowadays, significant progress has been made towards the goal of understanding the stability of SiCf/SiC composites under irradiation condition, including the fundamental properties such as strength and dimensional stability at high temperature [6,7]. However, the compatibility of SiC-based materials with coolant has been
* Corresponding author. Shanghai Institute of Ceramics, Chinese Academy of Sciences, No. 1295, Dingxi Road, Changning District, Shanghai 200050, China. ** Corresponding author. Shanghai Institute of Ceramics, Chinese Academy of Sciences, No. 1295, Dingxi Road, Changning District, Shanghai 200050, China. E-mail addresses: [email protected] (Z. Wang), [email protected] (S. Dong). http://dx.doi.org/10.1016/j.jnucmat.2017.02.007 0022-3115/© 2017 Published by Elsevier B.V.
demonstrated seldomly up to now, especially the database of corrosion behavior of SiCf/SiC composites in variety of coolants at high temperature. More recently, researchers [8e10] focus on the excellent compatibility of SiCf/SiC composites with molten fluoride salt (e.g. 46.5 mol % LiF- 11.5 mol % NaF-42.0 mol % KF eutectic molten salt, i.e. FLiNaK salt), one kind of candidate coolants used in the new Generation IV Nuclear Reactor [11,12]. This unique character make SiCf/SiC composites be one of the alternative materials for heat-exchanger, control rod and the other structure components in the new nuclear reactor concepts, such as the Molten Salt Reactor (MSR) and the Very High Temperature Reactor (VHTR) [13e15]. The corrosion of material in the liquid fluoride salt at high temperature is related to its composition. Free Si in C/SiSiC composite ceramic materials prepared by melt infiltration (MI) processing would be corroded selectively when exposed to the FLiNaK salt at 850 C for 500 h, while the SiC phase and remained pure C phase are relatively less attacked. This is qualitatively understood with the respect of the Gibbs free energy of fluoride formation per F2 molecule of the Si and SiC phases [16]. SiC fiber with higher oxygen content suffers worse corrosion in molten FLiNaK salt at
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700 C for 15 days than SiC fiber contained much fewer oxygen, indicating that the oxygen in SiC accelerates the corrosion of SiC in fluoride salt [10]. To meet the stricter service requirements in the highly corrosive liquid fluoride salt, the nuclear grade SiCf/SiC composites should be composed of high-purity, near-stoichiometric, and dense fibers and matrix. Various processes have been employed for SiCf/SiC composites fabrication, including chemical vapor infiltration (CVI), polymer impregnation and pyrolysis (PIP), reaction sintering (RS) and NanoInfiltration and Transient Eutectic phase (NITE) processing [17e20]. Among the various fabrication processes, the chemical vapor infiltration (CVI) processing is attractive for the fabrication of SiCf/ SiC composites because it yields highly crystalline, nearstoichiometric SiC matrix and the fiber damage during processing is minimized, which are widely employed to prepare the nuclear grade SiCf/SiC composites. However, the uniformity of deposited SiC matrix composition during CVI processing is determined by the temperature evenness in the furance, especially in the fabrication of the large-scale component. Owing to the existence of temperature gradient, the matrix located in the higher temperature region would be high-purity, near-stoichiometric SiC, and the matrix obtained in the lower temperature region would be composed of codeposited free Si and SiC [21]. The co-deposited free Si would affect the corrosion resistance performance of SiCf/SiC composites in liquid fluoride salt. Nevertheless, the effect of non-uniformity of matrix composition obtained in CVI processing to the corrosion behavior of SiCf/SiC composites has been report rarely so far, which is overwhelmingly important for the safe application of the material. In this work, SiCf/SiC composites with designed matrix, which is composed of Si-rich SiC layers and near-stoichiometric SiC layers, are prepared by controlling the CVI temperature to simulate the non-uniform SiC matrix caused by temperature gradient. The corrosion behavior of these designed SiC matrix in molten fluoride salt are studied. The microstructural evolution during the corrosion process, the corrosion kinetics at 800 C and the underlying corrosion processing are experimentally and theoretically investigated.
scale component. The fiber volume fraction and porosity of asreceived composites were ~40% and ~15%. The composites were cut into 35(long) 4(wide) 2(thick) mm3 bar for the corrosion experiment. Each corrosion experiment group contains 5 bars. The experimental setup of the corrosion test is schematically shown in Fig. 1. Due to the good chemical compatibility of graphite with fluoride molten salt, the corrosion test was carried out in a graphite crucible. Before corrosion, the graphite crucible and SiCf/ SiC samples were heat-treated in furnace at 700 C for 10 h to eliminate possible adsorptions. Therein, SiCf/SiC samples were fixed on a graphite support and dipped into ~80 g 46.5LiF - 11.5NaF - 42.0KF (mol. %, FLiNaK) fluoride eutectic molten salt (Shanghai Institute of Organic Chemistry, Chinese Academy of Sciences, Shanghai), which contained ~120 ppm O content. Subsequently, the crucibles were furnace heated at 800 C for 50 h, 100 h, 200 h, 300 h or 500 h, respectively. Here, it should be mentioned that all the above operations were conducted in a glove box under the protection of high purity argon gas. The atmosphere in the glove box was controlled at O2 <0.1 ppm and moisture <0.1 ppm during the corrosion process. After the high temperature corrosion process in fluoride eutectic molten salt, the SiCf/SiC composites samples were picked out from the crucible and immersed into 1 mol/L Al(NO3)3 solution to remove the fluoride salt solid adhering on the exterior and interior surfaces of the sample [23]. Finally, the samples were cleaned by ultrasonicating in deionized water, then dried in a vacuum drying chamber. The samples were weighed by precision balance (BT-25-S, 0.01 mg, Sartorius) before and after the corrosion test, the mass loss ratio, and mass loss rate, were applied to characterize the corrosion resistance property of the composites and corrosion kinetics. Phase compositions of the samples were analyzed by X-ray diffraction (XRD, D8 Advance, Bruker AXS Co. Ltd.) at room temperature, using CuKa radiation (l ¼ 0.154 nm) at a scanning 2q angles range of 10 e80 . The microstructure and chemical composition of samples before and after the corrosion test were characterized by scanning electron microscope (SEM, Quanta-250, FEI) and energy dispersive spectrometer (EDS, Oxford), respectively. 3. Results
2. Experimental The materials evaluated are two-dimensional woven fabrics of SiC fibers reinforced chemical vapor infiltrated SiC matrix composites. The used SiC fibers, known as KD-II SiC fibers, were provided by the National University of Defense Technology, Changsha, Hunan Province, China. The basic properties are shown in Table 1. More key characteristics of these SiC fibers, including the corrosion resistance of fibers in FLiNaK molten salt, have been reported elsewhere [10,22]. In advance of the CVI process, a pyrolytic carbon (PyC) monolayer about 100 nm in thickness was deposited on the SiC fiber surface through chemical vapor deposition (CVD). Using the Methyltrichlorosilane (MTS) as the raw material and H2 as the carrier gas, the SiCf/SiC composites with designed Si-rich SiC layers and near-stoichiometric SiC layers in matrix were prepared by controlling the deposited temperature via CVI processing. The purpose of preparation Si-rich SiC layers in SiC matrix was to simulate the non-uniformity of SiC matrix composition caused by temperature gradient in the CVI process for preparation the large
Fig. 2 shows mass loss ratio and mass loss rate of the samples after corrosion at 800 C for vary times in the molten FLiNaK salt. The mass loss ratio increased from 0.716 wt. % to 5.914 wt. % as the corrosion time was prolonged from 50 to 500 h, indicating a cumulative corrosion damage of the sample. The fitting line showed that the samples' mass loss ratio, ml, is as a function of the corrosion time, t, at 800 C. Linear relationship ml ¼ f(t) was observed when t was among 0e300 h. However, the mass loss ratio of samples after corrosion for 500 h deviated from the fitting line seriously. Furthermore, with the extended corrosion exposure, the mass loss rate showed a trend of first decrease and then increase and could be divided into three stages: the decelerated (I), constant rate (II) and accelerated (III) stages. The surface morphologies of the SiCf/SiC samples before and after corrosion are shown in Fig. 3. Compared with the flat surface of samples before corrosion, the sample surface became obviously rougher after corrosion, and matrix damage could be distinctively observed. For the sample corroded for 100 h, parts of the matrix
Table 1 General properties of the as-received KD-II SiC fibers [22]. Properties
C/Si mole ratio
Oxygen (wt %)
Density (g/cm3)
Diameter (mm)
Strength (GPa)
Modulus (GPa)
KD-II
1.41
1.3
2.7
12.1
3.0
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Fig. 1. Schematic illustration of (a) the experimental assembly and (b) the graphite crucible setup used for molten salt corrosion test.
Fig. 2. Mass loss ratio and mass loss rate of samples corroded at 800 C for various time. The minus sign of mass loss rate represents mass decrement.
layer on its surface were exfoliated as indicated by the arrows in Fig. 3(b). When the corrosion time was prolonged to 500 h, more intensive matrix corrosion and exfoliation occurred, which was in good agreement with the large increase in the mass loss ratio of the composites. At this time, SiC fibers with remnant thin SiC deposition layer exposed on the sample surface, and layer-exfoliation of the SiC layers were frequently observed (indicated by black arrow in Fig. 3). The cross section images of the samples before and after corrosion are shown in Fig. 4. Matrix with multilayer structure was observed in samples before and after corrosion. For the sample before corrosion, two types of designed matrix layers were found in the composites, the light gray matrix layer and deep gray matrix
layer. These matrix layers were separated with interlayer boundaries (as shown in white box in Fig. 4(a)). EDS analysis results shown in Table 2 revealed that the deep gray matrix layer showed atomic ratio of Si/C z 1.0 (point 1 in Fig. 4(a)), while the light gray one was significantly Si-rich (point 2 in Fig. 4(a)). This indicated that the deep gray matrix layer should be mainly composed of SiC, whereas a large amount of free Si should exist in the designed light gray matrix layer. Linescan analysis of interlayer boundary area in deposited matrix shown in the white box in Fig. 4(a) suggested stronger intensity of oxygen element in the boundary than matrix. In another word, the interlayer boundaries contained more oxygen than deposited matrix. After being exposed the fluoride molten salt at 800 C for 100 h (Fig. 4(b)), the light gray matrix layers and interlayer boundaries nearly disappeared, due to the molten salt corrosion. This suggested the poorer corrosion resistance of the Si-rich SiC layers and the interlayer boundaries in comparison with the SiC layers. EDS analysis of the corrosion regions (point 3 and 4 in Fig. 4(c)) revealed significant silicon losses, which made the chemical composition of these regions shift from silicon rich to carbon rich. Further corrosion of the composites led to the formation of vesicular structured layers in the matrix as the corrosion time was prolonged to 500 h, which contributed to the significant increment of the mass loss ratio. On another aspect, the oxygen content in corrosion area got higher than that in deposited layers of virgin samples. Fig. 5 shows the XRD patterns of the SiCf/SiC composites before and after corrosion at 800 C for 100 and 500 h. The samples before and after corrosion for 100 h showed quite similar XRD patterns, where diffraction peaks solely assigned to b-SiC were detected. While for the sample after corrosion for 500 h, new diffraction peak corresponding to Graphite-2H was identified in its XRD patter in addition to the peaks of b-SiC. Considering the high C content at the corrosion area obtained by EDS analysis, Graphite-2H may be the
Fig. 3. SEM images of the surface morphology of samples (a) before corrosion, and after corrosion at 800 C for (b) 100 h and (c) 500 h.
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Fig. 4. SEM images of the cross sections of the samples: (a) before corrosion, (b) corroded for 100 h, (c) higher magnification of area in white box in (b), and (d) corroded for 500 h.
Table 2 EDS Analysis of the local chemical composition on cross sections of the SiCf/SiC ceramic matrix composites (at. %). Point
1
2
3a
4a
C Si O
53.02 45.42 1.56
29.44 70.56 e
60.85 30.48 5.85
61.50 28.63 8.00
a
The F, K etc., elements of remained fluoride molten salt, with low content were not listed in the table.
corrosion product of b-SiC in the molten salt. 4. Discussion 4.1. Corrosion kinetics Solid reaction consists of many simple physical and chemical processes, such as chemical reaction, diffusion, crystallization, and so on [24]. As shown in Fig. 6(a), the solid reaction forms a reaction
Fig. 5. XRD patterns of the SiCf/SiC composites (a) before and (b) after corrosion at 800 C for 100 and 500 h.
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Fig. 6. Schematic drawing of solid corrosion models.
layer on the contact surface, for example, SiO2 layer formed by passive oxidation of Si or SiC [25e27]. The kinetics of solid reaction would be controlled by the slowest process, especially the reaction process and diffusion process [28]. For the corrosion of Si and SiC in liquid fluoride salt, the kinetics could be expressed by the equation below [24]:
n¼
1 kcn0
1 þ Dcx 0
(1)
where n is corrosion rate, c0 is the concentration of corrosion media, k is the corrosion reaction constant, n is the order of reaction, D is the diffusion constant of corrosion agent, and x is the thickness of formed reaction layer. When SiC or Si is dipped into high temperature molten fluoride salt, SiF4 is believed to be the corrosion production, which would dissolve into the liquid fluoride salt [29], and no dense layer would be obtained on the surface of SiCf/SiC composites in liquid fluoride salt (shown in Fig. 6(b)). As a result, there is no diffusion process during corrosion (x ¼ 0, D ¼ 0). In another word, the corrosion kinetics is controlled by reaction process
n ¼ kcn0
(2)
Moreover, the graphite, one of the corrosion production showed in XRD patterns, would have no contribution to the mass loss, which suggests that the source of mass loss would be the loss of Si from both free Si and SiC. The relationship between mass loss, Dm, and corrosion time, t, could be expressed as follows:
Dm ¼ MSi $
Z
ndt ¼ MSi $
Z kcn0 dt ¼ MSi kcn0 t
(3)
with Dm being mass loss of samples, t the corrosion time, and MSi the molar mass of Si. Then the relationship between mass loss ratio, ml, and corrosion time, t, could be obtained:
ml ¼
Dm m
¼
MSi k$cn0 $t m
(4)
where ml and m are the mass loss ratio and the original mass of samples, respectively. In consideration of the corrosion in pure liquid fluoride salt, the concentration of corrosion media, c0, would be invariant. That is to say, the corrosion kinetics of SiCf/SiC composites in liquid fluoride salt concides with linear law. This is in consist with the liner relationship between mass loss ratio, ml, and the corrosion time, t, when t is ranged from 0 to 300 h. However, the mass loss rate, which corresponds to the corrosion rate that should be constant, shows the trend of decrease (Stage I in Fig. 2) first, and then being stable from 100 h to 300 h in Stage II. The higher mass loss rate at first is associated with the quick dissolution of oxidation layer on surface of SiCf/SiC composites and the high O-content boundaries between deposited matrix layers, which leads to larger mass loss in short time. As the depletion of the O-contained compositions, the corrosion reaction was shifted to the direct corrosion of matrix with
much lower rate. As a result, the mass loss rate would be dominated by the simple matrix corrosion with relatively stable corrosion rate (Stage II). On the other aspect, the mass loss ratio of the samples corroded for 500 h deviated from the fitting line seriously, and the mass loss rate increases slightly (Stage III). This would be associated to the worse damage after corrosion in FLiNaK salt for 500 h. As shown in Fig. 3(c), damaged fibers with no SiC deposition layer in the samples corroded for 500 h exposed because of the exfoliation of deposition matrix layers. This may be imputed to that liquid FLiNaK salt would erode the Si-rich SiC layer and boundaries between the deposition layers at first. After being destroyed the Si-rich layers and boundaries, the bonding force between deposition layers is gone, which leads to the more severe layers exfoliation and higher mass loss than that caused by corrosion reaction only. 4.2. Microstructural evolution The microstructural evolution of the designed CVI matrix before and after corrosion in liquid LiNaKF salt at 800 C shown in Fig. 4 clearly demonstrates that there are two kinds of corrosion mechanism during the erosion of designed matrix. These two kinds of mechanism are totally different. In view of the non-uniformity of deposited matrix caused by temperature gradient in the CVI furnace [30], especially in the large-scale CVI instrument for large-scale components preparation, Si-rich SiC layers are designed and introduced into the matrix besides near-stoichiometric SiC layers by controlling CVI parameters [31] to simulate the non-uniform CVI SiC matrix. The microstructure of SiCf/SiC composites before and after corrosion at 800 C shown in Fig. 4 indicated that these two kinds of designed matrix layers have different corrosion behaviors. The designed Si-rich SiC layers suffer worse corrosion than near-stoichiometric SiC layers. This is corresponding to the research of Jens Schmidt et al. [16] conducted at 850 C. As reported in this research, a qualitative characteristic of the corrosion potential of the phases made by the Gibb's free energy of fluoride formation per F2 molecule (DGf) of the Si and SiC phases; the more negative the DGf for a certain phase, the more this phase can be easily corroded. Calculation results show that the DGf of SiF4 formed from Si (727 kJ/mol) is more negative than from SiC (696 kJ/mol), suggesting that the Si has poor corrosion resistance than SiC at 850 C. This result implies that the co-deposited Si in matrix is the weak area and should be avoided to the maximum extent during CVI processing for preparing the nuclear grade SiCf/SiC composites. During the preparation of SiC matrix by CVI processing, repeated grinding of the surface as it seal off and reimpregnating are usually needed for decreasing the porosity and obtaining denser composites [32]. Before the next infiltration processing, the deposited layer surface obtained prior infiltration would be oxidized slowly. When subsequent infiltration is finished, the very thin oxidation layer is covered by another deposited SiC layer and forms the higheroxygen-content boundary between deposited SiC matrix layers. During corrosion, the higher oxygen content deteriorates the corrosion resistance performance of boundaries, which is identical
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Fig. 7. Schematic diagrams showing the high temperature corrosion process of SiCf/SiC composites in fluoride molten salt.
to the results reported by Yang et al. [10]. Silicon oxide would be dissolved by fluoride, and the fluorine ion would attack the Si-O-Si bond and form Si-F bond [33], which leads to the selectively corrosion of layer boundaries. These corroded boundaries will be used as the penetration channels of liquid fluoride salt. As to the corrosion of SiC in fluoride molten salt, C seemed to be one of the corrosion products. The appearance of C maybe caused by the corrosion of SiC by either FLiNaK salt or the impurity, i.e. Fe and Ni, in fluoride salt, which needs to be studied further. 4.3. Corrosion mechanism Based on the above results, the corrosion process of the CVI SiC matrix of SiCf/SiC composites in high temperature fluoride molten salt can be schematically described as Fig. 7. First, the molten salt infiltrates into the sample through its surface defects, such as pores or cracks. The oxided surface and highe O-content interlayer boundaries are dissolved quickly. Therein, the Si-rich SiC matrix that may be introduced by temperature gradient will be preferentially corroded and will lead to the formation of fissure-like channels that allow the molten salt penetrating deeply into the matrix. After extended corrosion exposure, the Si-rich SiC layers are seriously corroded. Also, the bonding between deposited layers disappears due to the corrosion of boundaries between layers and the damage of Si-rich SiC layers. These corrosion damages result in the formation of pores in the matrix layers and the exfoliation of matrix layers around fibers. 5. Conclusions The high temperature corrosion behavior and microstructural evolution of SiCf/SiC composites with designed matrix in LiF-NaFKF eutectic molten salt at 800 C for vary corrosion time was studied. According to the experimental results, the following conclusions could be drawn: 1. The designed Si-rich SiC layers in matrix and O-contained boundaries between deposited layers are the weak area of the composites in terms of the corrosion resistance to fluoride molten salt. They are more liable to be corroded by the fluoride molten salt in comparison with the near-stoichiometric SiC layers. 2. The mass loss ratio induced by the molten salt corrosion shows an obvious increase with the extending of exposure time. The liner corrosion kinetics from 0 to 300 h is obtained during the corrosion of Si and SiC in FLiNaK at high temperature and is coincide with the reaction-control process. The deviation of 500 h samples is caused by the both reaction corrosion and SiC matrix layer exfoliation. 3. During the corrosion of the CVI SiC matrix in fluoride molten salt, the Si-rich SiC layers and O-contained boundaries of deposited layers are preferentially corroded, followed by the
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