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Corrosion behavior of ZrC–SiC composite ceramics in LiF–NaF–KF molten salt at high temperatures
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CERAMICS INTERNATIONAL
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Corrosion behavior of ZrC–SiC composite ceramics in LiF–NaF–KF molten salt at high temperatures Yang Wanga, Zhong-Feng Tanga,n, Yuan Fua, Shi-Rong Huangb,nn, Su-Fang Zhaoa, Peng Zhanga, Lei-Dong Xiea, Xin-Gang Wangc,nnn, Guo-Jun Zhangc a
Shanghai Institute of Applied Physics, Chinese Academy of Sciences, Shanghai 201800, China b College of Chemical Engineering, Xiangtan University, Xiangtan 411105, China c State Key Laboratory of High Performance Ceramics and Superfine Microstructures, Shanghai Institute of Ceramics, Chinese Academy of Sciences, Shanghai 200050, China Received 26 March 2015; received in revised form 16 June 2015; accepted 30 June 2015
Abstract The suitability of ZrC–SiC-based ceramics was assessed as candidate materials for fluoride. An immersion corrosion with ZrC–SiC composites was performed in molten FLiNaK salts (LiF–NaF–KF: 46.5/11.5/42 mol%) at 850 1C. Results indicated that ZrC–SiC-based ceramics were more corrosion resistant than the ZrC ceramic. ZrC–40SiC ceramic exhibited 50% weight-loss to ZrC. This carbide exhibited minimal corrosion attack because of strong resistance to dissolution in molten fluorides. Furthermore, XRD results were consistent with weight-loss and SEM results. The corrosion behavior of ZrC–SiC-based ceramics was affected by impurities, particularly oxygen content. Oxygen and/or oxygen ion participated in corrosion of ceramic. Corrosion of carbide ceramics is mainly due to C4 in ceramics. This ion is oxidized to elemental C by oxidants or impurities in the molten salt. The addition of C may inhibit corrosion of molten salts and thus improve the corrosion resistance of carbide ceramics. & 2015 Elsevier Ltd and Techna Group S.r.l. All rights reserved.
Keywords: C. Corrosion; Silicon carbide; Zirconium carbide; Molten FLiNaK salts; Micromorphology
1. Introduction Zirconium carbide (ZrC) possesses high melting temperature, high fracture strength, high electrical and thermal conductivities, and resistance to erosion/corrosion [1]. ZrC is a candidate material for reactors because of its excellent neutronic properties [2]. However, the poor chemical stability of ZrC at high temperatures in oxidizing atmosphere significantly limits its actual application [1,2]. A common approach to improve oxidation resistance is by incorporating Si-containing compounds into the ZrC matrix to form a protective SiO2n
Corresponding author. Tel./fax: þ 86 21 39195081. Corresponding author. nnn Corresponding author. E-mail addresses: [email protected] (Z.-F. Tang), [email protected] (S.-R. Huang), [email protected] (X.-G. Wang). nn
containing oxide scale [3–5]. Zhao et al. prepared composites comprising ZrC and 30 vol% silicon carbide (SiC) [4]. The protective effect of the oxide scales is enhanced by the formed SiO2 [4]. ZrC–SiC composites can generate high-performance ceramics; these composites demonstrate the passivating characteristic of SiC and the high melting temperature, hardness, and thermal stability of ZrC [3–5]. ZrC–SiC composites are broadly applied under extreme conditions and can be potentially use in fluoride salts because of their good thermomechanical properties, irradiation resistance, and chemical stability at high temperatures [2,6,7]. ZrC–SiC composites are suitable candidate materials in molten salt environments. The molten LiF–KF–NaF eutectic salt (46.5/11.5/42 mol%), commonly referred as FLiNaK, is a leading candidate material for heat transfer fluids because of its superior heat transfer properties [8,9]. FLiNaK molten salts can be used as a secondary coolant and a primary simulation coolant in molten
http://dx.doi.org/10.1016/j.ceramint.2015.06.143 0272-8842/& 2015 Elsevier Ltd and Techna Group S.r.l. All rights reserved.
Please cite this article as: Y. Wang, et al., Corrosion behavior of ZrC–SiC composite ceramics in LiF–NaF–KF molten salt at high temperatures, Ceramics International (2015), http://dx.doi.org/10.1016/j.ceramint.2015.06.143
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salt reactors (MSRs) [10,11]. FLiNaK molten salts can also be used as a primary coolant in advanced high-temperature reactors [12,13]. FLiNaK salts are used in MSRs because of their high thermal conductivities, high specific heats, low viscosities, and high boiling points. However, FLiNaK molten salts can be corrosive to materials at high temperatures; corrosion resistance under high-temperature environments is compromised by the thermodynamic instability of the protective oxide layer [14–16]. Ceramics exhibit excellent properties, such as high melting point, hardness, strength, and chemical stability. These attractive characteristics result in broad application of ceramics under extreme conditions. Ceramics are unique candidate materials that can be used at high temperatures (approximately 1000 1C) for a long time and satisfy the requirements of Generation IV nuclear reactors [17]. Related research work has been performed but cannot support the requirements to be classified as Generation IV nuclear reactors. Nishimura dipped SiC in an LiF–BeF2 (66/34 mol%) solution with traces of HF at 550 1C; after 10 days, a thick deposit with a depth of 500 nm was obtained, which exhibits very good fluoride corrosion resistance [18]. Schmidt and Peterson immersed SiC in FLiNaK molten salt at 850 1C to improve its corrosion resistance [19]. However, limited work has been conducted with regard to the corrosion of ZrC and ZrC–SiC composites in FLiNaK molten fluorides. The compatibility of fresh molten salts with materials is a technology gap for heat transfer applications; such materials derive their corrosion resistance from the formation of a protective surface oxide film, which is chemically unstable in most high-temperature fluoride salts. MSR materials are severely challenged by high operating temperatures and corrosive environments of molten fluoride salts; in particular, these materials are corroded in high-temperature molten fluorides. In this study, the static corrosion behavior of ZrC and ZrC– SiC-based ceramics dipped in FLiNaK molten salts is determined. The corrosion behavior and compatibility of ZrC–SiCbased ceramics with the FLiNaK molten salt mixture were elucidated. With thermodynamics as the basis, microstructural analysis and X-ray diffraction (XRD) were employed to explain the possible corrosion mechanism. Some interesting phenomena were reported which have not been found in previous works.
through a 200 mesh. The obtained mixed powders were hot pressed at 2100 1C. A pressure of 30 MPa was applied at 1450 1C to produce a rectangular sample with dimensions of 37 mm (length) 30 mm (width) 5 mm (height). The atmosphere was vacuum (o 10 Pa) at 1000 1C and then switched to flowing argon (99.99% pure). A detailed description of the experimental materials and processes can be found in the literature [1]. The bulk densities of sintered ceramics were measured using the Archimedes method. Final relative densities were determined as the ratio of the experimental bulk densities to the theoretical densities calculated using the rule of mixtures. The final relative densities of ZrC, ZrC–20SiC, ZrC– 40SiC, and SiC were 99.0%, 99.2%, 98.8%, and 98.6%, respectively. FLiNaK salt was purified with NH4HF at 750 1C under vacuum. In this purification process, residual water and oxygen were removed from the melt and metal impurities were minimized. The salt sample was analyzed using inductively coupled plasma–optical emission spectrometer (ICP–OES) before detecting the following impurities: 11 ppm Ca, 10 ppm Ni, 7 ppm Mg, and 101 ppm Si. Other elements were determined below 5 ppm or below the quantitative detection limits of ICP–OES. The impurities present in FLiNaK can significantly affect the corrosion behavior. For example, water present in the initial salt can significantly induce corrosion by producing HF. In the present study, the effects of trace impurities on corrosion were not specifically investigated to compare the relative corrosion performance of various materials. The concentrations of residual oxygen and water in FLiNaK were 1150 and 150 ppm as determined through a LECO oxygen analyzer and Karl–Fischer titration, respectively. The eutectic was then crushed into powder and stored in a glove box with argon gas. Test coupons were nominally cut into 12.5 mm (length) 10.0 mm (width) 3.0 mm (height) by using a wire-cut electrical discharge machine. The coupons were progressively ground with a 320-grit SiC sand paper and then with 1200-grit SiC. The coupons were finally polished using 1-μm diamond paste. After mechanical grinding, the coupons were degreased in an ultrasonic bath of acetone and alcohol for 5 min and then rinsed in deionized water before testing. 2.2. Corrosion procedure
2. Materials and methods 2.1. Materials ZrC (4 99% purity, 1.0 μm) and SiC ( 499% purity, 0.45 μm, Changle Xinyuan Carborundum Co., Ltd., China) starting powders were used; ZrC was synthesized through carbothermal reduction as described in a previous work [1,5]. ZrC, ZrC with different SiC contents (20 and 40 vol%), and SiC with 2.0 wt% B4C and 1.0 wt% C additives were mixed with ethanol in a plastic bottle through ball mixing of Si3N4 for 24 h; the corresponding samples were designated as ZrC, ZrC– 20SiC, ZrC–40SiC, and SiC. A rotary evaporator was used to remove ethanol at 70 1C. The powder mixtures were sieved
A schematic drawing of the experimental apparatus is shown in Fig. 1. The following procedures were performed in this apparatus. Three samples of a given test ceramic were fixed at each groove of graphite rods. The rods were placed in their corresponding graphite crucibles and then baked at 750 1C in an Ar environment for 8 h. All samples were tested in triplicates, with one crucible assigned to each ceramic. A schematic illustration of the design is shown in Fig. 1(b). The crucibles were filled with 50.0 g of FLiNaK salt before transferring to a furnace, where they were heated and maintained at 850 1C. After 100 h of exposure, the furnace temperature was decreased to room temperature and the FLiNaK salt was removed from the experimental apparatus.
Please cite this article as: Y. Wang, et al., Corrosion behavior of ZrC–SiC composite ceramics in LiF–NaF–KF molten salt at high temperatures, Ceramics International (2015), http://dx.doi.org/10.1016/j.ceramint.2015.06.143
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Fig. 1. Corrosion test apparatus.
The retrieved samples were cleaned of the residual FLiNaK by using 1 M Al(NO3)3 according to the method used by Kirillov and Fedulov [15]. After cleaning, the test coupons were baked in a vacuum oven for 24 h. The dimensions and weight of each coupon were measured and recorded. 2.3. Characterization Weight loss of the specimens before and after the corrosion test was measured using an electro-reading balance with an accuracy of 0.1 mg. Weight loss is defined by the following equation: Δm ¼ ðma mb Þ=S where ma is the mass of the sample after immersion, mb is the mass of the sample before immersion, and S is the surface area of the sample. Phases of the samples were investigated using XRD. XRD patterns were obtained at ambient temperature with a Rigaku D/max-RB diffractometer using Cu Kα radiation (λ ¼ 1.54 Å). Accelerating voltage and electric current were 36 kV and 25 mA, respectively. The samples were scanned from 101 to 801 with a step length of 0.021. Surface morphology was observed by SEM (Hitachi S5600LV, Japan) equipped with an Oxford energy-dispersive X-ray spectrometry (EDS). A cross section was cut at the span wise center, polished, and then subjected to cross-sectional analysis through SEM/EDS. 3. Results and discussion 3.1. Corrosion behavior studies The changes in weight of ceramics corroded in molten FLiNaK salt at 850 1C for 100 h are shown in Fig. 2. The weight loss of ZrC, ZrC–20SiC, ZrC–40SiC, and SiC caused by corrosion was of negative values. ZrC exhibited the maximum weight loss of about 100.0 mg/cm2. The weight loss of ZrC–20SiC, ZrC–40SiC, and SiC decreased to 95.5, 55.5, and 6.7 mg/cm2, respectively. The corrosion resistance of pristine ZrC was poor, whereas that of the ZrC–20SiC, ZrC– 40SiC, and SiC samples gradually increased in molten FLiNaK salt at 850 1C for 100 h. The ZrC–40SiC sample presented the highest corrosion resistance with a 50% weight loss to ZrC.
Fig. 2. Weight loss of ceramic samples caused by corrosion after exposure to FLiNaK at 850 1C for 100 h under Ar atmosphere.
These data confirmed that the ZrC and ZrC–20SiC samples demonstrated evident corrosion, cracking, and peeling, whereas the ZrC–40SiC and SiC samples remained intact relatively. Moreover, ZrC–40SiC exhibited more cracks and spalling on the surface and the SiC sample maintained a good specular gloss. Fig. 3 presents the images of the samples before and after corrosion testing. The first row refers to the pristine samples. All sample surfaces were smooth with specular gloss. The sample images after corrosion is presented in the second row. The ZrC and ZrC–20SiC samples demonstrated evident corrosion, cracking, and peeling. The ZrC–40SiC and SiC samples remained intact; the former presented more cracks and spalling on the surface, and the latter maintained a good specular gloss. The corrosive surface of SiC was similar to the surface before corrosion. 3.2. Crystalline structure change The XRD patterns of virgin and corroded samples after cleaning are presented in Fig. 4. The XRD patterns of ZrC are shown in Fig. 4(a). Five strong peaks were separately observed at 2θ ¼ 33.033 (1 1 1), 38.328 (2 0 0), 55.323 (2 2 0), 65.965 (3 1 1), and 69.304 (2 2 2). The observed peaks are in agreement with those of ZrC (XRD PDF Card no. 65-8408)/ (XRD PDF Card no. 35-0784). In the XRD patterns, no other
Please cite this article as: Y. Wang, et al., Corrosion behavior of ZrC–SiC composite ceramics in LiF–NaF–KF molten salt at high temperatures, Ceramics International (2015), http://dx.doi.org/10.1016/j.ceramint.2015.06.143
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Fig. 3. Images of different samples before (subscribe 1) and after (subscribe 2) corrosion testing in molten FLiNaK salt at 850 1C for 100 h: (a) ZrC, (b) ZrC– 20SiC, (c) ZrC–40SiC, and (d) SiC.
Fig. 4. XRD patterns of different samples before and after corrosion testing in molten FLiNaK salt at 850 1C for 100 h: (a) ZrC, (b) ZrC–20SiC, (c) ZrC–40SiC, and (d) SiC. Please cite this article as: Y. Wang, et al., Corrosion behavior of ZrC–SiC composite ceramics in LiF–NaF–KF molten salt at high temperatures, Ceramics International (2015), http://dx.doi.org/10.1016/j.ceramint.2015.06.143
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crystalline phase was found in the substrate before corrosion; hence, ZrC was in a very pure form. After corrosion, the peak at 26.603 identified C as a new phase. The peaks observed at 28, 30, 31 and 50 may be ascribed to ZrO2, and the other peaks corresponded to LiF, NaF, and KF. XRD analysis of ZrC– 20SiC is presented in Fig. 4(b), which demonstrated that the composites comprised ZrC and SiC. The dominant phase in the corrosion scale was SiC, and a strong peak assigned to C was detected at 26.603 (0 0 2). The peak intensity that corresponded to residual ZrC severely decreased. The low peak intensity of ZrC was also detected in the XRD spectrum. Fig. 4 (c) shows the XRD pattern collected on the washed surface of ZrC–40SiC before and after corrosion at 850 1C for 100 h. The main peaks in the XRD pattern were attributed to SiC, and the other weak peaks were due to C and FLiNaK. Moreover, the peaks of the ZrC substrate disappeared. The XRD pattern of corroded ZrC–40SiC was similar to that of SiC after corrosion. In this case, the substrate was selectively corroded and a corrosion scale thicker than the depth of the penetrating X-ray beam was formed. However, the peak ascribed to ZrO2 was not found in this pattern. The pattern of SiC is shown in Fig. 4(d). All strong intensity peaks before corrosion can be indexed to the 6H–SiC and 3C–SiC phases; these peaks are consistent with the values reported in the literature (XRD JCPDS Card no. 65-0962). No other phases of Si and C were further detected. The 6H–SiC phase did not evidently transform into 3C–SiC as indicated by the XRD pattern after corrosion [Fig. 4(d)]. Except the peaks of LiF, NaF, and KF and accompanied with traces of C, no other new phase that corresponded to the corrosion product was found in the substrate after corrosion testing for 100 h. This finding indicated the good chemical compatibility of SiC with molten FLiNaK salt. The XRD analyses showed that no other crystalline phase was found in the substrate before corrosion; hence, ZrC was very pure. After corrosion, a low amount of the C and ZrO2 phases could be detected. Phase composition strongly depended on the corrosion conditions. The EDS spectrum of ZrC–20SiC demonstrated that the composites comprised ZrC and SiC. The dominant phase after corrosion testing was SiC; however, the peaks that corresponded to the residual ZrC severely decreased. These results suggested that the thick C phase was formed on the specimen surface. The main peaks in the XRD pattern were obtained from SiC, the other weak peaks were attributed to C and FLiNaK, and the peaks for the ZrC substrate disappeared. The XRD pattern of corroded ZrC– 40SiC was similar to that of SiC after corrosion. In this case, the substrate was selectively corroded and a corrosion scale thicker than the depth of the penetrating X-ray beam was formed. This result was supported by EDX analysis on the samples exposed for 100 h. However, the peaks ascribed to ZrO2 were not detected in this pattern. All strong intensity peaks before corrosion could be indexed to the 6H–SiC and 3C–SiC phases and were in agreement with the value reported in the literature (JCPDS Card no. 65-0962). No other Si and C phases were observed. Except the peaks of LiF, NaF, and KF and accompanied with small traces of C, no other new phase
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corresponding to the corrosion product was found in the substrate after corrosion for 100 h, indicating the good chemical compatibility of SiC with molten FLiNaK salt. Furthermore, the experimental results were in good agreement with weight loss and SEM results.
3.3. Microstructural characterization The surface morphologies of the ceramics before and after corrosion testing in molten FLiNaK salt at 850 1C for 100 h are shown in Fig. 5. The surface morphology of the ZrC sample demonstrated grains with evident boundaries before corrosion. This surface morphology became considerably rougher after exposure to the molten salt and demonstrated a deep crack with a width of about 7 μm and a cavity. The resulting surface morphology indicated that ZrC underwent dissolution and etching. Fig. 5(b) shows the SEM images of the surface of ZrC–20SiC before and after corrosion testing. SiC content was presented as a dispersed phase in the ZrC continuous phase before corrosion testing and was similar to a sea island structure. After immersion in FLiNaK for 100 h, the structure became a sea–sea type, in which ZrC was no longer in the continuous phase but dispersed as particles in the porous SiC continuous phase; an approximately 2-μm crack was also found on the surface of ZrC. Fig. 5(c) shows that ZrC–40SiC exhibited a sea–sea structure, in which ZrC and SiC were presented as continuous phases that were uniformly distributed on the smooth surface. The porous SiC framework was retained by removing ZrC after exposure to FLiNak. Very few ZrC were residual. These findings indicated that ZrC was selectively attacked by molten FLiNaK salts. A surface morphology similar to ZrC was also observed in SiC before corrosion as shown in Fig. 5(d). The corrosion surface remained plain, and the cavity minimally increased. With the exception of the minor attack at the grain boundaries, this carbide exhibited minimal corrosion attack because of its strong resistance to dissolution in the molten fluoride salts. The fracture surface morphologies of different samples after corrosion testing in molten FLiNaK salt at 850 1C for 100 h are shown in Fig. 6. ZrC exhibited severe pitting corrosion, with more pits located near the surface. These pits presented diameters ranging from few to hundred micrometers and a corrosion depth of 1.5 mm in the entire cross section. The local impurity condensation in static FLiNaK and/or on the substrate causes local electric circuit, which results in pitting corrosion [20]. The ZrC–20SiC profile demonstrated different colors with an evident corrosion region as shown in Fig. 6(b). The depth of the external corrosion region was about 300 μm. Fig. 6(c) presents the profile of ZrC–40SiC after corrosion; this profile was similar to that of ZrC–20SiC and presented two phases with different colors. The corrosion region was about 150 μm. Furthermore, a row of corrosion notch with a depth of 30 μm was found on the edge of the cross-sectional SiC [Fig. 6(d)]. We inferred that adding SiC into ZrC can significantly reduce the depth of corrosion region and improve corrosion resistance. Corrosion resistance decreased in the
Please cite this article as: Y. Wang, et al., Corrosion behavior of ZrC–SiC composite ceramics in LiF–NaF–KF molten salt at high temperatures, Ceramics International (2015), http://dx.doi.org/10.1016/j.ceramint.2015.06.143
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Fig. 5. Surface morphology of different samples before (subscript 1) and after (subscript 2) corrosion testing in molten FLiNaK salt at 850 1C for 100 h: (a) ZrC, (b) ZrC–20SiC, (c) ZrC–40SiC, and (d) SiC.
order of ZrC o ZrC–20SiC oZrC–40SiC o SiC, which was consistent to the weight changes. The corrosive surface of SiC was similar to the surface before corrosion. The surface morphology of ZrC–20SiC was considerably rougher after exposure to molten FLiNak salts.
The surface morphology indicated that ZrC underwent dissolution and etching. The SiC content was presented as a dispersed phase in the ZrC continuous phase before corrosion testing and was similar to a sea island structure. During immersion in FLiNaK for 100 h, the structure was transformed
Please cite this article as: Y. Wang, et al., Corrosion behavior of ZrC–SiC composite ceramics in LiF–NaF–KF molten salt at high temperatures, Ceramics International (2015), http://dx.doi.org/10.1016/j.ceramint.2015.06.143
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Fig. 6. Fracture surface morphology of different samples after corrosion testing in molten FLiNaK salt at 850 1C for 100 h: (a) ZrC, (b) ZrC–20SiC, (c) ZrC–40SiC, and (d) SiC.
into a sea–sea type. ZrC was no longer presented as a continuous phase but dispersed as particles in the porous SiC continuous phase; an approximately 2-μm crack was also found on the surface. ZrC–40SiC exhibited a sea–sea structure, in which ZrC and SiC were presented as continuous phases that were uniformly distributed on the smooth surface. A porous SiC framework was retained by removing ZrC after exposure to FLiNaK, and very few ZrC were residual. These findings indicated that ZrC was selectively attacked by molten FLiNaK salts. A surface morphology similar to ZrC was further observed in SiC. The corrosion surface remained plain, and the cavity minimally increased. With the exception of minor attacks at the grain boundaries, this carbide exhibited minimal corrosion attack because of its strong resistance to dissolution in the molten salt fluorides. Hence, the corrosion of FLiNaK was influenced by impurities, particularly O contents. The concentration of O in FLiNaK salts after corrosion test was lower than that before the test. The O content of FLiNaK salt after immersing SiC, ZrC–40SiC, ZrC–20SiC, and ZrC samples decreased, and the corrosion resistance also decreased. This finding indicated that O and/or oxygen ion participated in corrosion of samples and crucible. The O in the molten salt was partially released as NO2 and as oxide adhered in the ceramics. After the test, the specimen surface showed a crevice-like state as shown in Figs.5 and 6. This state was possibly formed by corrosion at grain, block, and packet boundaries. The width of the remaining portion was almost similar to the block width [16,18]. These results were in agreement with the corrosion and SEM results. In addition, the surface was partially covered by an oxide layer. The resulting
high oxygen potential was attributed to the high concentration of H2O in FLiNaK. H2O dissolved in FLiNaK was utilized to form oxide layers. The formation mechanism of the oxide layer was possibly based on the typical oxide layer formation in high-temperature fluids as reported in previous study [16]. The EDS elemental distribution maps collected on the cross section of ZrC–40SiC corroded at 850 1C in the molten salt are presented in Fig. 7. The morphology of the Zr-depleted section [Fig. 7(a1)] was similar to that of pure SiC after exposure to FLiNaK. Moreover, the morphology of the non-corroded section [Fig. 7(a2)] was similar to that of ZrC–40SiC. EDS mapping performed on the cross-sectional images of ZrC– 40SiC showed that Zr was absent in the corrosion region and uniformly depleted in the transition region [Fig. 7(b)]. This finding confirmed that ZrC was selectively attacked by molten FLiNaK salt and removed or peeled from the substrate. K element penetrated into the substrate by about 150 μm as shown in Fig. 7(c). This finding was consistent with the Zr depletion depth on BSE and EDX maps of SEM. Corrosion attack was particularly drastic in the ZrC phase, whereas the SiC phase was relatively less attacked by the molten salts. This result was also confirmed through point element analysis of the cross section of ZrC and SiC after corrosion. Fig. 8(a) shows the point element of SiC after corrosion. Few element components differed between inner corrosion and outer non-corrosion, with Si:C near to 1:1. This finding implied that SiC was not corroded. However, element components of ZrC evidently changed before and after immersion [Fig. 8(b)]; in particular (in Zr and O), Zr:C diverged far from the 1:1 stoichiometry. The corrosion
Please cite this article as: Y. Wang, et al., Corrosion behavior of ZrC–SiC composite ceramics in LiF–NaF–KF molten salt at high temperatures, Ceramics International (2015), http://dx.doi.org/10.1016/j.ceramint.2015.06.143
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Fig. 7. Cross-section micrographs, EDS elemental line scans, and the corresponding elemental distribution maps of ZrC–40SiC after corrosion testing in molten FLiNaK salt at 850 1C for 100 h. (a)SiC and (b)ZrC.
potential of the phases can be qualitatively analyzed by considering Gibb's free energy of oxidation product formation per oxidant molecule (△G) of the ZrC and SiC phases [16]. A negative △G value for a particular phase could result in more phases that were prone to attack. At 850 1C, the △G for ZrO2 was more negative when formed from pure ZrC ( 349.1 kJ) than SiF4 when formed from SiC ( 347.8 kJ); hence, the ZrC phase was more prone to dissolution than the SiC phase. Gibb's free energy of formation data was calculated using the Facsage computer software program and its associated databases. According to the electronegativity formula reported by Pauling, molten salt is a type of ionic liquid; the iconicity of the Si–C bond (12%) was lower than that of the Zr–C bond (31%); hence, iconicity may also affect the corrosion resistance of molten salts. Therefore, high contents of ZrC in ceramic materials can result in severe corrosion. 3.4. Corrosion mechanisms Fig. 9 shows that O content changed in FLiNaK salts after corrosion with different ceramic materials. The O content of the FLiNaK salt after immersing SiC, ZrC–40SiC, ZrC–20SiC, and ZrC samples decreased, and the corrosion resistance also decreased. The nitrogen content comes from the FLiNaK salt due to a preparation method. This finding indicated that the corrosion of FLiNaK was affected by impurities, particularly O content. Hence, O and/or oxygen ion participate in corrosion of samples and crucible. A negative correlation existed between the oxygen content of salt and material corrosion. O in molten FLiNaK salt partially escaped as NO2 and as oxide adhered in
ceramics. This result indicated that the O (such as NO3 and H2O) content in molten FLiNaK salt participated in corrosion of ceramic samples. We can infer that corrosion may be mainly caused by the reaction of ZrC with oxidizing substances or impurities, particularly NO3 . NO3 was oxidized with carbide to form C and oxide on the interface to release NO2 (g) from the molten salt. Few SiC was also simultaneously oxidized, and the residual or hydrolyzed HF also corroded the carbides. The oxide reacted with HF or dissolved by molten FLiNaK salt. Element line scan of the ZrC–40SiC composite after corrosion is shown in Fig. 7(a). The C content in the corrosive area significantly increased, whereas Zr content significantly decreased. Moreover, O, F, K, and Si contents minimally increased. The combined element distribution of the corroded section [Fig. 7(b) and (c)] and XRD analysis (Fig. 4) is also presented. The corrosion mechanism of ZrC–SiC composite in molten FLiNaK salt at 850 1C is shown in Fig. 10. The corrosion of ZrC–SiC composite ceramics in molten fluorides is mainly due to the negative valence of C4 in ceramic materials; these ions were oxidized to elemental C by oxidizing substances or impurities in the molten salt. The generated fluoride finally coordinated with alkaline fluoride molten salts to form [ZrF7]2 and [SiF6]2 . The Si content in the profiles of the ceramics increased because ZrC was selectively corroded, whereas SiC was retained; hence, the Si content relatively increased. The corrosive behavior of ceramic composites was similar to that of graphite. The sample fix was significantly lower than that of graphite and maintained gloss luster, indicating that C can inhibit corrosion. This finding indirectly
Please cite this article as: Y. Wang, et al., Corrosion behavior of ZrC–SiC composite ceramics in LiF–NaF–KF molten salt at high temperatures, Ceramics International (2015), http://dx.doi.org/10.1016/j.ceramint.2015.06.143
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Fig. 8. Point element analysis of the cross section of ZrC and SiC after corrosion testing in molten FLiNaK salt at 850 1C for 100 h: (a)SiC and (b)ZrC.
Fig. 9. Oxygen content of the FLiNaK salts changed after immersion of different ceramics.
confirmed that the C4 in ceramics reduced by oxidation substances or contaminant in molten salt could be attributed for the molten salt corrosion of the ceramic material. We can infer that adding C to carbide ceramic materials should enhance their corrosion resistance in molten salts. 4. Conclusion An immersion corrosion test using ZrC–SiC-based ceramics was performed in molten FLiNaK salts at 850 1C for 100 h by using a graphite-resistance furnace connected with a glove box to ensure salt purity. ZrC–SiC-based ceramics exhibited higher corrosion resistance than ZrC ceramics. ZrC–40SiC ceramics demonstrated a 50% weight loss to ZrC ceramics. The ZrC and ZrC–20SiC samples presented evident corrosion, cracking, and
Fig. 10. Corrosion mechanism of ZrC–SiC-based composite in molten FLiNaK salt at 850 1C.
peeling, whereas the ZrC–40SiC and SiC samples remained intact and maintained a good specular gloss. The surface morphology of ZrC–20SiC was considerably rougher after exposure to molten FLiNak salts. ZrC–40SiC exhibited a sea– sea structure, in which ZrC and SiC were presented as continuous phases that were uniformly distributed on the smooth surface. XRD results suggested that the thick C phase was formed on the surface of ZrC–SiC-based specimens. Furthermore, XRD experimental results were in a good agreement with weight loss and SEM results. The corrosion behavior of ZrC–SiC-based ceramics was affected by FLiNaK salt impurities, particularly oxygen contents. Hence, oxygen and/or oxygen ions participated in corrosion of the ceramic
Please cite this article as: Y. Wang, et al., Corrosion behavior of ZrC–SiC composite ceramics in LiF–NaF–KF molten salt at high temperatures, Ceramics International (2015), http://dx.doi.org/10.1016/j.ceramint.2015.06.143
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samples. In ZrC–SiC-based composites, ZrC was selectively attacked by molten salt and corrosion resistance improved with increasing SiC content. The corrosion mechanism of carbide ceramics was mainly attributed to C4 in ceramic materials; these ions were oxidized to elemental C by oxidizing substances or impurities in the molten salt. Acknowledgments The project was supported by the “Strategic Priority Research Program” of the Chinese Academy of Sciences, Grant no. XD02002400. Financial supports from the National Natural Science Foundation of China (Nos. 21406256 and 11205229). Key Laboratory of Nuclear Radiation and Nuclear Energy Technology, Chinese Academy of Sciences and the State Key Laboratory of High Performance Ceramics and Superfine Microstructures are also greatly appreciated. References [1] X.G. Wang, J.X. Liu, Y.M. Kan, G.J. Zhang, Effect of solid solution formation on densification of hot-pressed ZrC ceramics with MC (M ¼V, Nb, and Ta) additions, J. Eur. Ceram. Soc. 32 (2012) 1795–1802. [2] C.R.F. Azevedo, Selection of fuel cladding material for nuclear fission reactors, Eng. Fail. Anal. 18 (2011) 1943–1962. [3] L. Zhao, D. Jia, X. Duan, Z. Yang, Y. Zhou, Low temperature sintering of ZrC–SiC composite, J. Alloy. Compd. 509 (2011) 9816–9820. [4] L. Zhao, D. Jia, X. Duan, Z. Yang, Y. Zhou, Oxidation of ZrC-30 vol% SiC composite in air from low to ultrahigh temperature, J. Eur. Ceram. Soc. 32 (2012) 947–954. [5] X.G. Wang, G.J. Zhang, J.X. Xue, Y. Tang, X. Huang, C.M. Xu, P.L. Wang, Reactive hot pressing of ZrC–SiC ceramics at low temperature, J. Am. Ceram. Soc. 96 (2013) 32–36. [6] S. Ueta, J. Aihara, K. Sawa, A. Yasuda, M. Honda, N. Furihata, Development of high temperature gas-cooled reactor (HTGR) fuel in Japan, Prog. Nucl. Energy 53 (2011) 788–793. [7] D. Pizon, R. Lucas, S. Chehaidi, S. Foucaud, A. Maître, From trimethylvinylsilane to ZrC–SiC hybrid materials, J. Eur. Ceram. Soc. 31 (2011) 2687–2690.
[8] D.E. Holcomb, S.M. Cetiner, An overview of liquid fluoride salt heat transport systems. ORNL/TM-2010/156 (2010) 1–70. [9] J. Ambrosek, M. Anderson, K. Sridharan, T. Allen, Current status of knowledge of the fluoride salt (FLiNaK) heat transfer, Nucl. Technol. 165 (2009) 166–173. [10] V. Khokhlov, V. Ignatiev, V. Afonichkin, Evaluating physical properties of molten salt reactor fluoride mixtures, J. Fluor. Chem. 130 (2009) 30–37. [11] T. Goto, T. Nohira, R. Hagiwara, Y. Ito, Selected topics of molten fluorides in the field of nuclear engineering, J. Fluor. Chem. 130 (2009) 102–107. [12] C. Forsberg, The advanced high-temperature reactor: high-temperature fuel, liquid salt coolant, liquid-metal-reactor plant, Prog. Nucl. Energy 47 (2005) 32–43. [13] M. Salanne, C. Simon, P. Turq, P.A. Madden, Heat-transport properties of molten fluorides: determination from first-principles, J. Fluor. Chem. 130 (2009) 38–44. [14] D.T. Ingersoll, C.W. Forsberg, L.J. Ott, D.F. Williams, J.P. Renier, D.F. Wilson, S.J. Ball, L. Reid, W.R. Corwin, G.D. Del Cul, P.F. Peterson, H. Zhao, P.S. Pickard, E.J. Parma, M. Vernon, Status of preconceptual design of the advanced high-temperature reactor (AHTR). ORNL/TM-2004/104 (2004) 1–86. [15] V.B. Kirillov, V.I. Fedulov, The corrosion resistance of 12Kh18N10T steel in molten fluoride salts, Mater. Sci. 16 (1981) 22–25. [16] M. Kondo, T. Nagasaka, V. Tsisar, A. Sagara, T. Muroga, T. Watanabe, T. Oshima, Y. Yokoyama, H. Miyamoto, E. Nakamura, N. Fujii, Corrosion of reduced activation ferritic martensitic steel JLF-1 in purified Flinak at static and flowing conditions, Fusion Eng. Des. 85 (2010) 1430–1436. [17] L.C. Olson, J.W. Ambrosek, K. Sridharan, M.H. Anderson, T.R. Allen, Materials corrosion in molten LiF–NaF–KF salt, J. Fluor. Chem. 130 (2009) 67–73. [18] H. Nishimura, T. Terai, T. Yoneoka, S. Tanaka, A. Sagara, O. Motojima, Compatibility of structural candidate materials with LiF–BeF2 molten salt mixture, J. Nucl. Mater. 283–287 (2000) 1326–1331. [19] J. Schmidt, M. Scheiffele, M. Crippa, P.F. Peterson., E. Urquiza, K. Sridharan, L.C. Olson, M.H. Anderson, T.R. Allen, Y. Chen, Design, fabrication, and testing of ceramic plate-type heat exchangers with integrated flow channel design, Int. J. Appl. Ceram. Technol. 8 (2011) 1073–1086. [20] F.Y. Ouyang, C.H. Chang, B.C. You, T.K. Yeh, J.J. Kai, Effect of moisture on corrosion of Ni-based alloys in molten alkali fluoride FLiNaK salt environments, J. Nucl. Mater. 437 (2013) 201–207.
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CERAMICS INTERNATIONAL
Ceramics International ] (]]]]) ]]]–]]] www.elsevier.com/locate/ceramint
Corrosion behavior of ZrC–SiC composite ceramics in LiF–NaF–KF molten salt at high temperatures Yang Wanga, Zhong-Feng Tanga,n, Yuan Fua, Shi-Rong Huangb,nn, Su-Fang Zhaoa, Peng Zhanga, Lei-Dong Xiea, Xin-Gang Wangc,nnn, Guo-Jun Zhangc a
Shanghai Institute of Applied Physics, Chinese Academy of Sciences, Shanghai 201800, China b College of Chemical Engineering, Xiangtan University, Xiangtan 411105, China c State Key Laboratory of High Performance Ceramics and Superfine Microstructures, Shanghai Institute of Ceramics, Chinese Academy of Sciences, Shanghai 200050, China Received 26 March 2015; received in revised form 16 June 2015; accepted 30 June 2015
Abstract The suitability of ZrC–SiC-based ceramics was assessed as candidate materials for fluoride. An immersion corrosion with ZrC–SiC composites was performed in molten FLiNaK salts (LiF–NaF–KF: 46.5/11.5/42 mol%) at 850 1C. Results indicated that ZrC–SiC-based ceramics were more corrosion resistant than the ZrC ceramic. ZrC–40SiC ceramic exhibited 50% weight-loss to ZrC. This carbide exhibited minimal corrosion attack because of strong resistance to dissolution in molten fluorides. Furthermore, XRD results were consistent with weight-loss and SEM results. The corrosion behavior of ZrC–SiC-based ceramics was affected by impurities, particularly oxygen content. Oxygen and/or oxygen ion participated in corrosion of ceramic. Corrosion of carbide ceramics is mainly due to C4 in ceramics. This ion is oxidized to elemental C by oxidants or impurities in the molten salt. The addition of C may inhibit corrosion of molten salts and thus improve the corrosion resistance of carbide ceramics. & 2015 Elsevier Ltd and Techna Group S.r.l. All rights reserved.
Keywords: C. Corrosion; Silicon carbide; Zirconium carbide; Molten FLiNaK salts; Micromorphology
1. Introduction Zirconium carbide (ZrC) possesses high melting temperature, high fracture strength, high electrical and thermal conductivities, and resistance to erosion/corrosion [1]. ZrC is a candidate material for reactors because of its excellent neutronic properties [2]. However, the poor chemical stability of ZrC at high temperatures in oxidizing atmosphere significantly limits its actual application [1,2]. A common approach to improve oxidation resistance is by incorporating Si-containing compounds into the ZrC matrix to form a protective SiO2n
Corresponding author. Tel./fax: þ 86 21 39195081. Corresponding author. nnn Corresponding author. E-mail addresses: [email protected] (Z.-F. Tang), [email protected] (S.-R. Huang), [email protected] (X.-G. Wang). nn
containing oxide scale [3–5]. Zhao et al. prepared composites comprising ZrC and 30 vol% silicon carbide (SiC) [4]. The protective effect of the oxide scales is enhanced by the formed SiO2 [4]. ZrC–SiC composites can generate high-performance ceramics; these composites demonstrate the passivating characteristic of SiC and the high melting temperature, hardness, and thermal stability of ZrC [3–5]. ZrC–SiC composites are broadly applied under extreme conditions and can be potentially use in fluoride salts because of their good thermomechanical properties, irradiation resistance, and chemical stability at high temperatures [2,6,7]. ZrC–SiC composites are suitable candidate materials in molten salt environments. The molten LiF–KF–NaF eutectic salt (46.5/11.5/42 mol%), commonly referred as FLiNaK, is a leading candidate material for heat transfer fluids because of its superior heat transfer properties [8,9]. FLiNaK molten salts can be used as a secondary coolant and a primary simulation coolant in molten
http://dx.doi.org/10.1016/j.ceramint.2015.06.143 0272-8842/& 2015 Elsevier Ltd and Techna Group S.r.l. All rights reserved.
Please cite this article as: Y. Wang, et al., Corrosion behavior of ZrC–SiC composite ceramics in LiF–NaF–KF molten salt at high temperatures, Ceramics International (2015), http://dx.doi.org/10.1016/j.ceramint.2015.06.143
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salt reactors (MSRs) [10,11]. FLiNaK molten salts can also be used as a primary coolant in advanced high-temperature reactors [12,13]. FLiNaK salts are used in MSRs because of their high thermal conductivities, high specific heats, low viscosities, and high boiling points. However, FLiNaK molten salts can be corrosive to materials at high temperatures; corrosion resistance under high-temperature environments is compromised by the thermodynamic instability of the protective oxide layer [14–16]. Ceramics exhibit excellent properties, such as high melting point, hardness, strength, and chemical stability. These attractive characteristics result in broad application of ceramics under extreme conditions. Ceramics are unique candidate materials that can be used at high temperatures (approximately 1000 1C) for a long time and satisfy the requirements of Generation IV nuclear reactors [17]. Related research work has been performed but cannot support the requirements to be classified as Generation IV nuclear reactors. Nishimura dipped SiC in an LiF–BeF2 (66/34 mol%) solution with traces of HF at 550 1C; after 10 days, a thick deposit with a depth of 500 nm was obtained, which exhibits very good fluoride corrosion resistance [18]. Schmidt and Peterson immersed SiC in FLiNaK molten salt at 850 1C to improve its corrosion resistance [19]. However, limited work has been conducted with regard to the corrosion of ZrC and ZrC–SiC composites in FLiNaK molten fluorides. The compatibility of fresh molten salts with materials is a technology gap for heat transfer applications; such materials derive their corrosion resistance from the formation of a protective surface oxide film, which is chemically unstable in most high-temperature fluoride salts. MSR materials are severely challenged by high operating temperatures and corrosive environments of molten fluoride salts; in particular, these materials are corroded in high-temperature molten fluorides. In this study, the static corrosion behavior of ZrC and ZrC– SiC-based ceramics dipped in FLiNaK molten salts is determined. The corrosion behavior and compatibility of ZrC–SiCbased ceramics with the FLiNaK molten salt mixture were elucidated. With thermodynamics as the basis, microstructural analysis and X-ray diffraction (XRD) were employed to explain the possible corrosion mechanism. Some interesting phenomena were reported which have not been found in previous works.
through a 200 mesh. The obtained mixed powders were hot pressed at 2100 1C. A pressure of 30 MPa was applied at 1450 1C to produce a rectangular sample with dimensions of 37 mm (length) 30 mm (width) 5 mm (height). The atmosphere was vacuum (o 10 Pa) at 1000 1C and then switched to flowing argon (99.99% pure). A detailed description of the experimental materials and processes can be found in the literature [1]. The bulk densities of sintered ceramics were measured using the Archimedes method. Final relative densities were determined as the ratio of the experimental bulk densities to the theoretical densities calculated using the rule of mixtures. The final relative densities of ZrC, ZrC–20SiC, ZrC– 40SiC, and SiC were 99.0%, 99.2%, 98.8%, and 98.6%, respectively. FLiNaK salt was purified with NH4HF at 750 1C under vacuum. In this purification process, residual water and oxygen were removed from the melt and metal impurities were minimized. The salt sample was analyzed using inductively coupled plasma–optical emission spectrometer (ICP–OES) before detecting the following impurities: 11 ppm Ca, 10 ppm Ni, 7 ppm Mg, and 101 ppm Si. Other elements were determined below 5 ppm or below the quantitative detection limits of ICP–OES. The impurities present in FLiNaK can significantly affect the corrosion behavior. For example, water present in the initial salt can significantly induce corrosion by producing HF. In the present study, the effects of trace impurities on corrosion were not specifically investigated to compare the relative corrosion performance of various materials. The concentrations of residual oxygen and water in FLiNaK were 1150 and 150 ppm as determined through a LECO oxygen analyzer and Karl–Fischer titration, respectively. The eutectic was then crushed into powder and stored in a glove box with argon gas. Test coupons were nominally cut into 12.5 mm (length) 10.0 mm (width) 3.0 mm (height) by using a wire-cut electrical discharge machine. The coupons were progressively ground with a 320-grit SiC sand paper and then with 1200-grit SiC. The coupons were finally polished using 1-μm diamond paste. After mechanical grinding, the coupons were degreased in an ultrasonic bath of acetone and alcohol for 5 min and then rinsed in deionized water before testing. 2.2. Corrosion procedure
2. Materials and methods 2.1. Materials ZrC (4 99% purity, 1.0 μm) and SiC ( 499% purity, 0.45 μm, Changle Xinyuan Carborundum Co., Ltd., China) starting powders were used; ZrC was synthesized through carbothermal reduction as described in a previous work [1,5]. ZrC, ZrC with different SiC contents (20 and 40 vol%), and SiC with 2.0 wt% B4C and 1.0 wt% C additives were mixed with ethanol in a plastic bottle through ball mixing of Si3N4 for 24 h; the corresponding samples were designated as ZrC, ZrC– 20SiC, ZrC–40SiC, and SiC. A rotary evaporator was used to remove ethanol at 70 1C. The powder mixtures were sieved
A schematic drawing of the experimental apparatus is shown in Fig. 1. The following procedures were performed in this apparatus. Three samples of a given test ceramic were fixed at each groove of graphite rods. The rods were placed in their corresponding graphite crucibles and then baked at 750 1C in an Ar environment for 8 h. All samples were tested in triplicates, with one crucible assigned to each ceramic. A schematic illustration of the design is shown in Fig. 1(b). The crucibles were filled with 50.0 g of FLiNaK salt before transferring to a furnace, where they were heated and maintained at 850 1C. After 100 h of exposure, the furnace temperature was decreased to room temperature and the FLiNaK salt was removed from the experimental apparatus.
Please cite this article as: Y. Wang, et al., Corrosion behavior of ZrC–SiC composite ceramics in LiF–NaF–KF molten salt at high temperatures, Ceramics International (2015), http://dx.doi.org/10.1016/j.ceramint.2015.06.143
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Fig. 1. Corrosion test apparatus.
The retrieved samples were cleaned of the residual FLiNaK by using 1 M Al(NO3)3 according to the method used by Kirillov and Fedulov [15]. After cleaning, the test coupons were baked in a vacuum oven for 24 h. The dimensions and weight of each coupon were measured and recorded. 2.3. Characterization Weight loss of the specimens before and after the corrosion test was measured using an electro-reading balance with an accuracy of 0.1 mg. Weight loss is defined by the following equation: Δm ¼ ðma mb Þ=S where ma is the mass of the sample after immersion, mb is the mass of the sample before immersion, and S is the surface area of the sample. Phases of the samples were investigated using XRD. XRD patterns were obtained at ambient temperature with a Rigaku D/max-RB diffractometer using Cu Kα radiation (λ ¼ 1.54 Å). Accelerating voltage and electric current were 36 kV and 25 mA, respectively. The samples were scanned from 101 to 801 with a step length of 0.021. Surface morphology was observed by SEM (Hitachi S5600LV, Japan) equipped with an Oxford energy-dispersive X-ray spectrometry (EDS). A cross section was cut at the span wise center, polished, and then subjected to cross-sectional analysis through SEM/EDS. 3. Results and discussion 3.1. Corrosion behavior studies The changes in weight of ceramics corroded in molten FLiNaK salt at 850 1C for 100 h are shown in Fig. 2. The weight loss of ZrC, ZrC–20SiC, ZrC–40SiC, and SiC caused by corrosion was of negative values. ZrC exhibited the maximum weight loss of about 100.0 mg/cm2. The weight loss of ZrC–20SiC, ZrC–40SiC, and SiC decreased to 95.5, 55.5, and 6.7 mg/cm2, respectively. The corrosion resistance of pristine ZrC was poor, whereas that of the ZrC–20SiC, ZrC– 40SiC, and SiC samples gradually increased in molten FLiNaK salt at 850 1C for 100 h. The ZrC–40SiC sample presented the highest corrosion resistance with a 50% weight loss to ZrC.
Fig. 2. Weight loss of ceramic samples caused by corrosion after exposure to FLiNaK at 850 1C for 100 h under Ar atmosphere.
These data confirmed that the ZrC and ZrC–20SiC samples demonstrated evident corrosion, cracking, and peeling, whereas the ZrC–40SiC and SiC samples remained intact relatively. Moreover, ZrC–40SiC exhibited more cracks and spalling on the surface and the SiC sample maintained a good specular gloss. Fig. 3 presents the images of the samples before and after corrosion testing. The first row refers to the pristine samples. All sample surfaces were smooth with specular gloss. The sample images after corrosion is presented in the second row. The ZrC and ZrC–20SiC samples demonstrated evident corrosion, cracking, and peeling. The ZrC–40SiC and SiC samples remained intact; the former presented more cracks and spalling on the surface, and the latter maintained a good specular gloss. The corrosive surface of SiC was similar to the surface before corrosion. 3.2. Crystalline structure change The XRD patterns of virgin and corroded samples after cleaning are presented in Fig. 4. The XRD patterns of ZrC are shown in Fig. 4(a). Five strong peaks were separately observed at 2θ ¼ 33.033 (1 1 1), 38.328 (2 0 0), 55.323 (2 2 0), 65.965 (3 1 1), and 69.304 (2 2 2). The observed peaks are in agreement with those of ZrC (XRD PDF Card no. 65-8408)/ (XRD PDF Card no. 35-0784). In the XRD patterns, no other
Please cite this article as: Y. Wang, et al., Corrosion behavior of ZrC–SiC composite ceramics in LiF–NaF–KF molten salt at high temperatures, Ceramics International (2015), http://dx.doi.org/10.1016/j.ceramint.2015.06.143
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Fig. 3. Images of different samples before (subscribe 1) and after (subscribe 2) corrosion testing in molten FLiNaK salt at 850 1C for 100 h: (a) ZrC, (b) ZrC– 20SiC, (c) ZrC–40SiC, and (d) SiC.
Fig. 4. XRD patterns of different samples before and after corrosion testing in molten FLiNaK salt at 850 1C for 100 h: (a) ZrC, (b) ZrC–20SiC, (c) ZrC–40SiC, and (d) SiC. Please cite this article as: Y. Wang, et al., Corrosion behavior of ZrC–SiC composite ceramics in LiF–NaF–KF molten salt at high temperatures, Ceramics International (2015), http://dx.doi.org/10.1016/j.ceramint.2015.06.143
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crystalline phase was found in the substrate before corrosion; hence, ZrC was in a very pure form. After corrosion, the peak at 26.603 identified C as a new phase. The peaks observed at 28, 30, 31 and 50 may be ascribed to ZrO2, and the other peaks corresponded to LiF, NaF, and KF. XRD analysis of ZrC– 20SiC is presented in Fig. 4(b), which demonstrated that the composites comprised ZrC and SiC. The dominant phase in the corrosion scale was SiC, and a strong peak assigned to C was detected at 26.603 (0 0 2). The peak intensity that corresponded to residual ZrC severely decreased. The low peak intensity of ZrC was also detected in the XRD spectrum. Fig. 4 (c) shows the XRD pattern collected on the washed surface of ZrC–40SiC before and after corrosion at 850 1C for 100 h. The main peaks in the XRD pattern were attributed to SiC, and the other weak peaks were due to C and FLiNaK. Moreover, the peaks of the ZrC substrate disappeared. The XRD pattern of corroded ZrC–40SiC was similar to that of SiC after corrosion. In this case, the substrate was selectively corroded and a corrosion scale thicker than the depth of the penetrating X-ray beam was formed. However, the peak ascribed to ZrO2 was not found in this pattern. The pattern of SiC is shown in Fig. 4(d). All strong intensity peaks before corrosion can be indexed to the 6H–SiC and 3C–SiC phases; these peaks are consistent with the values reported in the literature (XRD JCPDS Card no. 65-0962). No other phases of Si and C were further detected. The 6H–SiC phase did not evidently transform into 3C–SiC as indicated by the XRD pattern after corrosion [Fig. 4(d)]. Except the peaks of LiF, NaF, and KF and accompanied with traces of C, no other new phase that corresponded to the corrosion product was found in the substrate after corrosion testing for 100 h. This finding indicated the good chemical compatibility of SiC with molten FLiNaK salt. The XRD analyses showed that no other crystalline phase was found in the substrate before corrosion; hence, ZrC was very pure. After corrosion, a low amount of the C and ZrO2 phases could be detected. Phase composition strongly depended on the corrosion conditions. The EDS spectrum of ZrC–20SiC demonstrated that the composites comprised ZrC and SiC. The dominant phase after corrosion testing was SiC; however, the peaks that corresponded to the residual ZrC severely decreased. These results suggested that the thick C phase was formed on the specimen surface. The main peaks in the XRD pattern were obtained from SiC, the other weak peaks were attributed to C and FLiNaK, and the peaks for the ZrC substrate disappeared. The XRD pattern of corroded ZrC– 40SiC was similar to that of SiC after corrosion. In this case, the substrate was selectively corroded and a corrosion scale thicker than the depth of the penetrating X-ray beam was formed. This result was supported by EDX analysis on the samples exposed for 100 h. However, the peaks ascribed to ZrO2 were not detected in this pattern. All strong intensity peaks before corrosion could be indexed to the 6H–SiC and 3C–SiC phases and were in agreement with the value reported in the literature (JCPDS Card no. 65-0962). No other Si and C phases were observed. Except the peaks of LiF, NaF, and KF and accompanied with small traces of C, no other new phase
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corresponding to the corrosion product was found in the substrate after corrosion for 100 h, indicating the good chemical compatibility of SiC with molten FLiNaK salt. Furthermore, the experimental results were in good agreement with weight loss and SEM results.
3.3. Microstructural characterization The surface morphologies of the ceramics before and after corrosion testing in molten FLiNaK salt at 850 1C for 100 h are shown in Fig. 5. The surface morphology of the ZrC sample demonstrated grains with evident boundaries before corrosion. This surface morphology became considerably rougher after exposure to the molten salt and demonstrated a deep crack with a width of about 7 μm and a cavity. The resulting surface morphology indicated that ZrC underwent dissolution and etching. Fig. 5(b) shows the SEM images of the surface of ZrC–20SiC before and after corrosion testing. SiC content was presented as a dispersed phase in the ZrC continuous phase before corrosion testing and was similar to a sea island structure. After immersion in FLiNaK for 100 h, the structure became a sea–sea type, in which ZrC was no longer in the continuous phase but dispersed as particles in the porous SiC continuous phase; an approximately 2-μm crack was also found on the surface of ZrC. Fig. 5(c) shows that ZrC–40SiC exhibited a sea–sea structure, in which ZrC and SiC were presented as continuous phases that were uniformly distributed on the smooth surface. The porous SiC framework was retained by removing ZrC after exposure to FLiNak. Very few ZrC were residual. These findings indicated that ZrC was selectively attacked by molten FLiNaK salts. A surface morphology similar to ZrC was also observed in SiC before corrosion as shown in Fig. 5(d). The corrosion surface remained plain, and the cavity minimally increased. With the exception of the minor attack at the grain boundaries, this carbide exhibited minimal corrosion attack because of its strong resistance to dissolution in the molten fluoride salts. The fracture surface morphologies of different samples after corrosion testing in molten FLiNaK salt at 850 1C for 100 h are shown in Fig. 6. ZrC exhibited severe pitting corrosion, with more pits located near the surface. These pits presented diameters ranging from few to hundred micrometers and a corrosion depth of 1.5 mm in the entire cross section. The local impurity condensation in static FLiNaK and/or on the substrate causes local electric circuit, which results in pitting corrosion [20]. The ZrC–20SiC profile demonstrated different colors with an evident corrosion region as shown in Fig. 6(b). The depth of the external corrosion region was about 300 μm. Fig. 6(c) presents the profile of ZrC–40SiC after corrosion; this profile was similar to that of ZrC–20SiC and presented two phases with different colors. The corrosion region was about 150 μm. Furthermore, a row of corrosion notch with a depth of 30 μm was found on the edge of the cross-sectional SiC [Fig. 6(d)]. We inferred that adding SiC into ZrC can significantly reduce the depth of corrosion region and improve corrosion resistance. Corrosion resistance decreased in the
Please cite this article as: Y. Wang, et al., Corrosion behavior of ZrC–SiC composite ceramics in LiF–NaF–KF molten salt at high temperatures, Ceramics International (2015), http://dx.doi.org/10.1016/j.ceramint.2015.06.143
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Fig. 5. Surface morphology of different samples before (subscript 1) and after (subscript 2) corrosion testing in molten FLiNaK salt at 850 1C for 100 h: (a) ZrC, (b) ZrC–20SiC, (c) ZrC–40SiC, and (d) SiC.
order of ZrC o ZrC–20SiC oZrC–40SiC o SiC, which was consistent to the weight changes. The corrosive surface of SiC was similar to the surface before corrosion. The surface morphology of ZrC–20SiC was considerably rougher after exposure to molten FLiNak salts.
The surface morphology indicated that ZrC underwent dissolution and etching. The SiC content was presented as a dispersed phase in the ZrC continuous phase before corrosion testing and was similar to a sea island structure. During immersion in FLiNaK for 100 h, the structure was transformed
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Fig. 6. Fracture surface morphology of different samples after corrosion testing in molten FLiNaK salt at 850 1C for 100 h: (a) ZrC, (b) ZrC–20SiC, (c) ZrC–40SiC, and (d) SiC.
into a sea–sea type. ZrC was no longer presented as a continuous phase but dispersed as particles in the porous SiC continuous phase; an approximately 2-μm crack was also found on the surface. ZrC–40SiC exhibited a sea–sea structure, in which ZrC and SiC were presented as continuous phases that were uniformly distributed on the smooth surface. A porous SiC framework was retained by removing ZrC after exposure to FLiNaK, and very few ZrC were residual. These findings indicated that ZrC was selectively attacked by molten FLiNaK salts. A surface morphology similar to ZrC was further observed in SiC. The corrosion surface remained plain, and the cavity minimally increased. With the exception of minor attacks at the grain boundaries, this carbide exhibited minimal corrosion attack because of its strong resistance to dissolution in the molten salt fluorides. Hence, the corrosion of FLiNaK was influenced by impurities, particularly O contents. The concentration of O in FLiNaK salts after corrosion test was lower than that before the test. The O content of FLiNaK salt after immersing SiC, ZrC–40SiC, ZrC–20SiC, and ZrC samples decreased, and the corrosion resistance also decreased. This finding indicated that O and/or oxygen ion participated in corrosion of samples and crucible. The O in the molten salt was partially released as NO2 and as oxide adhered in the ceramics. After the test, the specimen surface showed a crevice-like state as shown in Figs.5 and 6. This state was possibly formed by corrosion at grain, block, and packet boundaries. The width of the remaining portion was almost similar to the block width [16,18]. These results were in agreement with the corrosion and SEM results. In addition, the surface was partially covered by an oxide layer. The resulting
high oxygen potential was attributed to the high concentration of H2O in FLiNaK. H2O dissolved in FLiNaK was utilized to form oxide layers. The formation mechanism of the oxide layer was possibly based on the typical oxide layer formation in high-temperature fluids as reported in previous study [16]. The EDS elemental distribution maps collected on the cross section of ZrC–40SiC corroded at 850 1C in the molten salt are presented in Fig. 7. The morphology of the Zr-depleted section [Fig. 7(a1)] was similar to that of pure SiC after exposure to FLiNaK. Moreover, the morphology of the non-corroded section [Fig. 7(a2)] was similar to that of ZrC–40SiC. EDS mapping performed on the cross-sectional images of ZrC– 40SiC showed that Zr was absent in the corrosion region and uniformly depleted in the transition region [Fig. 7(b)]. This finding confirmed that ZrC was selectively attacked by molten FLiNaK salt and removed or peeled from the substrate. K element penetrated into the substrate by about 150 μm as shown in Fig. 7(c). This finding was consistent with the Zr depletion depth on BSE and EDX maps of SEM. Corrosion attack was particularly drastic in the ZrC phase, whereas the SiC phase was relatively less attacked by the molten salts. This result was also confirmed through point element analysis of the cross section of ZrC and SiC after corrosion. Fig. 8(a) shows the point element of SiC after corrosion. Few element components differed between inner corrosion and outer non-corrosion, with Si:C near to 1:1. This finding implied that SiC was not corroded. However, element components of ZrC evidently changed before and after immersion [Fig. 8(b)]; in particular (in Zr and O), Zr:C diverged far from the 1:1 stoichiometry. The corrosion
Please cite this article as: Y. Wang, et al., Corrosion behavior of ZrC–SiC composite ceramics in LiF–NaF–KF molten salt at high temperatures, Ceramics International (2015), http://dx.doi.org/10.1016/j.ceramint.2015.06.143
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Fig. 7. Cross-section micrographs, EDS elemental line scans, and the corresponding elemental distribution maps of ZrC–40SiC after corrosion testing in molten FLiNaK salt at 850 1C for 100 h. (a)SiC and (b)ZrC.
potential of the phases can be qualitatively analyzed by considering Gibb's free energy of oxidation product formation per oxidant molecule (△G) of the ZrC and SiC phases [16]. A negative △G value for a particular phase could result in more phases that were prone to attack. At 850 1C, the △G for ZrO2 was more negative when formed from pure ZrC ( 349.1 kJ) than SiF4 when formed from SiC ( 347.8 kJ); hence, the ZrC phase was more prone to dissolution than the SiC phase. Gibb's free energy of formation data was calculated using the Facsage computer software program and its associated databases. According to the electronegativity formula reported by Pauling, molten salt is a type of ionic liquid; the iconicity of the Si–C bond (12%) was lower than that of the Zr–C bond (31%); hence, iconicity may also affect the corrosion resistance of molten salts. Therefore, high contents of ZrC in ceramic materials can result in severe corrosion. 3.4. Corrosion mechanisms Fig. 9 shows that O content changed in FLiNaK salts after corrosion with different ceramic materials. The O content of the FLiNaK salt after immersing SiC, ZrC–40SiC, ZrC–20SiC, and ZrC samples decreased, and the corrosion resistance also decreased. The nitrogen content comes from the FLiNaK salt due to a preparation method. This finding indicated that the corrosion of FLiNaK was affected by impurities, particularly O content. Hence, O and/or oxygen ion participate in corrosion of samples and crucible. A negative correlation existed between the oxygen content of salt and material corrosion. O in molten FLiNaK salt partially escaped as NO2 and as oxide adhered in
ceramics. This result indicated that the O (such as NO3 and H2O) content in molten FLiNaK salt participated in corrosion of ceramic samples. We can infer that corrosion may be mainly caused by the reaction of ZrC with oxidizing substances or impurities, particularly NO3 . NO3 was oxidized with carbide to form C and oxide on the interface to release NO2 (g) from the molten salt. Few SiC was also simultaneously oxidized, and the residual or hydrolyzed HF also corroded the carbides. The oxide reacted with HF or dissolved by molten FLiNaK salt. Element line scan of the ZrC–40SiC composite after corrosion is shown in Fig. 7(a). The C content in the corrosive area significantly increased, whereas Zr content significantly decreased. Moreover, O, F, K, and Si contents minimally increased. The combined element distribution of the corroded section [Fig. 7(b) and (c)] and XRD analysis (Fig. 4) is also presented. The corrosion mechanism of ZrC–SiC composite in molten FLiNaK salt at 850 1C is shown in Fig. 10. The corrosion of ZrC–SiC composite ceramics in molten fluorides is mainly due to the negative valence of C4 in ceramic materials; these ions were oxidized to elemental C by oxidizing substances or impurities in the molten salt. The generated fluoride finally coordinated with alkaline fluoride molten salts to form [ZrF7]2 and [SiF6]2 . The Si content in the profiles of the ceramics increased because ZrC was selectively corroded, whereas SiC was retained; hence, the Si content relatively increased. The corrosive behavior of ceramic composites was similar to that of graphite. The sample fix was significantly lower than that of graphite and maintained gloss luster, indicating that C can inhibit corrosion. This finding indirectly
Please cite this article as: Y. Wang, et al., Corrosion behavior of ZrC–SiC composite ceramics in LiF–NaF–KF molten salt at high temperatures, Ceramics International (2015), http://dx.doi.org/10.1016/j.ceramint.2015.06.143
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Fig. 8. Point element analysis of the cross section of ZrC and SiC after corrosion testing in molten FLiNaK salt at 850 1C for 100 h: (a)SiC and (b)ZrC.
Fig. 9. Oxygen content of the FLiNaK salts changed after immersion of different ceramics.
confirmed that the C4 in ceramics reduced by oxidation substances or contaminant in molten salt could be attributed for the molten salt corrosion of the ceramic material. We can infer that adding C to carbide ceramic materials should enhance their corrosion resistance in molten salts. 4. Conclusion An immersion corrosion test using ZrC–SiC-based ceramics was performed in molten FLiNaK salts at 850 1C for 100 h by using a graphite-resistance furnace connected with a glove box to ensure salt purity. ZrC–SiC-based ceramics exhibited higher corrosion resistance than ZrC ceramics. ZrC–40SiC ceramics demonstrated a 50% weight loss to ZrC ceramics. The ZrC and ZrC–20SiC samples presented evident corrosion, cracking, and
Fig. 10. Corrosion mechanism of ZrC–SiC-based composite in molten FLiNaK salt at 850 1C.
peeling, whereas the ZrC–40SiC and SiC samples remained intact and maintained a good specular gloss. The surface morphology of ZrC–20SiC was considerably rougher after exposure to molten FLiNak salts. ZrC–40SiC exhibited a sea– sea structure, in which ZrC and SiC were presented as continuous phases that were uniformly distributed on the smooth surface. XRD results suggested that the thick C phase was formed on the surface of ZrC–SiC-based specimens. Furthermore, XRD experimental results were in a good agreement with weight loss and SEM results. The corrosion behavior of ZrC–SiC-based ceramics was affected by FLiNaK salt impurities, particularly oxygen contents. Hence, oxygen and/or oxygen ions participated in corrosion of the ceramic
Please cite this article as: Y. Wang, et al., Corrosion behavior of ZrC–SiC composite ceramics in LiF–NaF–KF molten salt at high temperatures, Ceramics International (2015), http://dx.doi.org/10.1016/j.ceramint.2015.06.143
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samples. In ZrC–SiC-based composites, ZrC was selectively attacked by molten salt and corrosion resistance improved with increasing SiC content. The corrosion mechanism of carbide ceramics was mainly attributed to C4 in ceramic materials; these ions were oxidized to elemental C by oxidizing substances or impurities in the molten salt. Acknowledgments The project was supported by the “Strategic Priority Research Program” of the Chinese Academy of Sciences, Grant no. XD02002400. Financial supports from the National Natural Science Foundation of China (Nos. 21406256 and 11205229). Key Laboratory of Nuclear Radiation and Nuclear Energy Technology, Chinese Academy of Sciences and the State Key Laboratory of High Performance Ceramics and Superfine Microstructures are also greatly appreciated. References [1] X.G. Wang, J.X. Liu, Y.M. Kan, G.J. Zhang, Effect of solid solution formation on densification of hot-pressed ZrC ceramics with MC (M ¼V, Nb, and Ta) additions, J. Eur. Ceram. Soc. 32 (2012) 1795–1802. [2] C.R.F. Azevedo, Selection of fuel cladding material for nuclear fission reactors, Eng. Fail. Anal. 18 (2011) 1943–1962. [3] L. Zhao, D. Jia, X. Duan, Z. Yang, Y. Zhou, Low temperature sintering of ZrC–SiC composite, J. Alloy. Compd. 509 (2011) 9816–9820. [4] L. Zhao, D. Jia, X. Duan, Z. Yang, Y. Zhou, Oxidation of ZrC-30 vol% SiC composite in air from low to ultrahigh temperature, J. Eur. Ceram. Soc. 32 (2012) 947–954. [5] X.G. Wang, G.J. Zhang, J.X. Xue, Y. Tang, X. Huang, C.M. Xu, P.L. Wang, Reactive hot pressing of ZrC–SiC ceramics at low temperature, J. Am. Ceram. Soc. 96 (2013) 32–36. [6] S. Ueta, J. Aihara, K. Sawa, A. Yasuda, M. Honda, N. Furihata, Development of high temperature gas-cooled reactor (HTGR) fuel in Japan, Prog. Nucl. Energy 53 (2011) 788–793. [7] D. Pizon, R. Lucas, S. Chehaidi, S. Foucaud, A. Maître, From trimethylvinylsilane to ZrC–SiC hybrid materials, J. Eur. Ceram. Soc. 31 (2011) 2687–2690.
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Please cite this article as: Y. Wang, et al., Corrosion behavior of ZrC–SiC composite ceramics in LiF–NaF–KF molten salt at high temperatures, Ceramics International (2015), http://dx.doi.org/10.1016/j.ceramint.2015.06.143