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Oxidation resistant SiC coating for graphite materials
Carbon 37 (1999) 1475–1484
Oxidation resistant SiC coating for graphite materials Qingshan Zhu*, Xueliang Qiu, Changwen Ma Institute of Nuclear Energy Technology, Tsinghua University, Beijing 100080, People’ s Republic of China Received 25 November 1998; accepted 8 January 1999
Abstract A dense functional gradient SiSiC coating with good oxidation resistance when applied to a graphite substrate has been developed using a novel process, in which large pores in the coating were eliminated by application of a primary coating, and pinholes were filled with free silicon. The mechanism for the formation of a continuous SiC coating on graphite was investigated experimentally and theoretically. Oxidation tests were conducted from 8008C to 16008C for up to 1000 h. Samples of graphite with the SiSiC coating were found to be totally intact after oxidizing for 1000 h together with 100 thermal cycles at 12008C and 200 h with 20 thermal cycles at 14008C, indicating excellent oxidation resistance and thermal shock resistance of the coating. The good thermal shock resistance is attributed to the compositional gradient in the coating. 1999 Elsevier Science Ltd. All rights reserved. Keywords: A. Graphite; B. Coating; Oxidation
1. Introduction Carbon and graphite are attractive materials for high temperature applications due to their high strength, high modulus, excellent thermal shock resistance and light weight. They are widely used as engineering materials, such as heaters, electrical contacts, high-temperature heat exchangers, rocket nozzles and leading edges of aircraft wings, etc. [1,2]. However, the use of carbon materials has been greatly restricted due to the poor oxidation resistance at high temperature in an oxidizing atmosphere. Achieving good oxidation resistance is crucial to utilize their full potential as high-temperature materials. Oxidation protection for carbon materials has been extensively studied in the past 60 years [3–5]. Ceramic coatings are commonly employed to protect carbon materials from oxidation [6–8]. Although several coating systems have been developed, SiC is considered to be the best coating material due to its good mechanical properties, low density, good physical–chemical compatibility with carbon and excellent oxidation resistance below 18008C [9]. The SiC coating can be formed by several methods [10–13]. Reaction-formed process, in which molten silicon reacts with the substrate on the surface to form an SiC coating, is one of the effective ways to form an SiC coating [14]. However, the SiC coating obtained from reaction-formed *Corresponding author.
process often contains many defects (pores, pinholes or cracks), and the oxidation resistance of the coating is not sufficient at high temperature, because oxygen can attack the substrate through these defects. Additional processes were usually needed to fill these defects. These processes are complicated and required longer time than the original reaction-formed process [15,16]. Another problem is that the coating formation behavior can be quite different when different carbon materials are used; sometimes it is difficult to form an SiC coating by this process [17]. Therefore, the coating formation mechanism still needs to be clarified. In the present work, the mechanism for the formation of the SiC coating was systematically studied. A new process, which was aimed at making a dense SiC coating directly, was developed and optimized. A dense SiSiC coating with a gradient SiC / C layer was formed on two kinds of graphite. The oxidation behavior of the coatings was tested from 8008C to 16008C for up to 1000 h.
2. Experimental
2.1. Materials Two kinds of graphite were used in this research. One was the matrix graphite of spherical fuel elements of a high temperature gas cooled reactor (HTGR), which is
0008-6223 / 99 / $ – see front matter 1999 Elsevier Science Ltd. All rights reserved. PII: S0008-6223( 99 )00010-X
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under construction at the authors’ institute. The matrix graphite is made of 64 wt.% natural flake graphite, 16 wt.% petroleum coke graphite and 20 wt.% resin binder [18]. This graphite will be termed ‘HTGR graphite’ hereafter. The other was a type of nuclear-grade graphite provided by Shanghai Carbon Corporation (Shanghai, P.R. China). It will be termed ‘Shanghai graphite’ in the rest of this paper. The characteristics of the graphite materials are summarized in Table 1. The graphite was cut into small specimens with a size of 2031538 mm. The surface of the specimens was polished with 5.0 mm diamond paste, followed by ultrasonically washing in acetone for 30 min. Other materials used in the process were Si powder (circa 1 mm), SiC powder (a, circa 0.1 mm) and nucleargrade graphite powder (circa 1 mm). Several reagents were also employed in the slip process. These reagents were (1) the acrylamide [CH 2 =CHCONH 2 ], the dispersant and monomer; (2) N,N9-methylene bisacrylamide [(CH 2 =CHCONH) 2 CH 2 ], the crosslinking agent; (3) ammonium persulfate [(NH 4 ) 2 S 2 O 8 ], the free-radical initiator for polymerization; and (4) tetramethylammonium hydroxide [(CH 3 ) 4 NOH], the agent for adjusting the pH value of the slip.
2.2. Coating preparation procedure A new process for preparing an oxidation resistant coating for graphite materials was developed from the conventional reaction-formed process [16] and the gelcasting process [19]. The preparation procedure of the process is shown in Fig. 1. The process consisted of three steps. First, a primary SiC coating was coated on the surface of a substrate, then the substrate with the primary coating was encased in a slip of silicon to form a packed gel, and finally, the packed body was sintered to produce the SiC coating. To produce a primary coating on a substrate, a-SiC powder, graphite powder and water were mixed with a specified amount of dispersant and crosslinking agent to form a slurry. The pH value of the slurry was adjusted to 10–12. The slip was milled for a sufficient time before the addition of the initiator. The slip was then sprayed on the surface of the substrate. The monomer in the slip polymerized to form a primary SiC coating, which could attach strongly to the substrate after drying. A packing slip, which consisted of Si powder, a-SiC
Table 1 Characteristics of graphite materials Graphite
Shanghai HTGR
r kg / m 3
1810 1730
Ash content 310 6
400 130–190
Thermal expansion coefficient 310 6 / K '
i
5.50 2.77
4.27 2.54
Fig. 1. Coating preparation procedure.
powder and water, was mixed and milled with a specified amount of dispersant and crosslinking agent. The substrate with primary coating was then dipped into the packing slip, and in the meantime, the initiator was added to the packing slip to initiate the polymerization reaction. A packed body, which contained the substrate, formed after the polymerization of the monomer. The speed of polymerization could be controlled by adjusting the temperature and the concentration of the initiator and crosslinking agent. The packed body was dried in an oven between 508C to 2008C. No special care is needed in the drying process. The dried body has very high strength, for example, it can even be machined if desired [19]. Consequently, there are no handling problems with the dried body. The oxidation resistant SiC coating could be obtained by sintering the packing body to a temperature between 15008C and 19508C for 1 to 4 h.
2.3. Coating characterization The surface of the coatings was observed by optical microscopy and scanning electron microscopy (SEM). Coating quality was assessed through examining the surface for cracks, pores and pinholes. Elemental distribution of Si in the cross section of the coatings was detected by electron probe microanalysis (EPMA) to study the distribution of the SiC phase. The crystalline phases of the coatings were evaluated by X-ray diffractometer (XRD). The density of the SiSiC coating, which was used to calculate the amount of free silicon in the coating through Eq. (1) by assuming that the coating was composed only of Si and SiC, was measured by a titration method [20].
Q. Zhu et al. / Carbon 37 (1999) 1475 – 1484
rm 5 rSi x 1 rSiC (1 2 x)
(1)
where rm is the measured density of the SiSiC coating, rSi the density of Si, rSiC the density of SiC and x the volume percentage of Si in the SiSiC coating. Before measuring the density, the SiSiC coating was burned at 6008C for a sufficient time to remove the attached graphite. The burned coating was ultrasonically washed in acetone for a sufficient time to remove the ash, followed by drying in an oven. Diiodomethane (CH 2 I 2 ) and bromoform (CHBr 3 ) were used as the titrants. The density of diiodomethane and bromoform is 3.326310 3 kg / m 3 , 2.893310 3 kg / m 3 and 3.320310 3 kg / m 3 , 2.889310 3 kg / m 3 at 48C and 208C, respectively.
2.4. Oxidation tests Thermogravimetric tests were carried out under natural convection conditions in air between 8008C and 10008C. The sample was supported by a silica basket hung on a silica chain from a Mettler balance. The sensitivity of the balance is 0.1 mg. The mass change was continuously recorded by a PC computer. Thermal cyclic tests at high temperature were carried out in a MoSi 2 furnace. The furnace was first heated to a specified temperature (12008C, 14008C and 16008C), then a sample was placed into the furnace and held for a specified time (10 h for the tests at 12008C and 14008C, and 5 h for 16008C). After oxidation, the sample was taken out and cooled at room temperature. The surface temperature of the sample can fall down to 2008C within 5 min. After cooling to room temperature, the sample was weighed and placed in the furnace again.
3. Results and discussion
3.1. Coating formation mechanism The reaction mechanism between solid carbon and molten silicon has been studied by various authors [21– 23]. Although SiC was proved to be formed through a solution–reprecipitation mechanism on a micro-scale, different results can be generated in forming the SiC coating by the reaction. For example, both Minnear [24] and Yamamoto et al. [15] reported that a continuous SiC layer could be formed by this reaction, while Hase and Suzuki [17] found no continuous SiC layer between the carbon particles and silicon and suggested that the reaction product quickly spalled due to the volume misfit between SiC and the carbon [22,25]. In our experiments, the authors also found that a continuous SiC coating was easy to form on Shanghai graphite, while it was difficult to form a continuous SiC coating on HGTR graphite by this reaction. In the case of HTGR graphite, big cracks and severe spallation occurred. Sometimes the substrate was crushed
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into powder by the reaction. The volume will expand |234% [26] when the graphite is totally converted into SiC. This volume expansion is believed to be the reason for the cracks and spallation, but the volume misfit of Shanghai graphite should be larger than that of HTGR graphite, because the density of Shanghai graphite is greater than that of HTGR graphite. Various investigations have been conducted to find the direct reason for this. It was found that the competition between the infiltration process and the chemical reaction process determined whether a continuous SiC coating was formed. If the infiltration rate is smaller than that of the reaction rate, the whole process is controlled by the infiltration process. The infiltrant (molten silicon) will be consumed at the infiltration frontier. During infiltration, the reaction-induced volume expansion will reduce the paths of infiltration, which further decreases the infiltration rate. Finally, the paths of infiltration will be totally blocked by the volume expansion and a continuous SiC coating forms. After that, the growth of the SiC layer is controlled by the diffusion of Si and C through the newly formed SiC layer. In this case, the reaction is limited only to the surface of the substrate. Therefore, the total volume expansion is small. If, on the other hand, the infiltration rate is much higher than the chemical reaction rate, the infiltrant can be delivered rapidly ahead of the reaction front. A large amount of molten silicon will penetrate deeply into the substrate before the infiltration paths are narrowed by reaction-induced volume expansion. The substrate will be converted into SiC not only on the surface but also inside the substrate, which results in a big volume expansion. This big volume expansion will cause cracks or spallation of the substrate even if a layer of SiC has already formed on the surface of the substrate. Cracks may also be caused by the thermal stress resulting from rapid local heating due to the quick reaction. The cracks can serve as the new paths for infiltration. More infiltrant can penetrate into the substrate through the newly formed cracks, and a big volume expansion can be generated by the consequent reaction, which may cause new cracks again. Such cycles can continue until the substrate is completely converted into SiC. The chemical reaction rate is a function of temperature and the graphite. The infiltration speed v can be expressed by the following equation [27]: dl rgl n cosu n 5 ] 5 ]]] dt 4hl
(2)
where t is the infiltration time, l the length of infiltration at time t, r the equivalent radius of the pore, glv the surface tension, u the wetting angle and h the viscosity of the infiltrant. When the system and experimental conditions are fixed, the chemical reaction rate is fixed, and glv , u and h are also fixed. The rate of infiltration is proportional to the pore diameter as indicated by Eq. (2). Therefore, the
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competition of these two processes is mainly determined by the pore size of the substrate. This point was confirmed by the experimental results of Singh and Behrendt [28], which revealed that the speed of infiltration and reaction was almost the same in the case of a small pore size (0.2 mm), while in the case of a medium pore size (1.8 mm), the infiltration speed was much higher than the reaction rate. Our investigations showed that the difference between Shanghai graphite and HTGR graphite in coating formation mainly lay in the pore structure, because the chemical kinetics of the two kinds of graphite with molten silicon was almost the same, as confirmed by differential thermal analyses (DTA) [29]. The pore size distribution of the two kinds of graphite is shown in Fig. 2. The pore size of Shanghai graphite is mainly over the range of 0.3–0.9 mm, which is very close to the small pore size of 0.2 mm examined by Singh and Behrendt [28]. In this case, the speed of the infiltration and the reaction must be in the same order, and a continuous SiC coating is expected to form. As seen in Fig. 2, the pore size of HTGR graphite shows a bimodal distribution, and the pores are mainly in the range of 2–4 mm, which is larger than the medium pore size indicated by Singh and Behrendt [28]. In this case, the infiltration speed must be much higher than the reaction rate. A continuous SiC coating is difficult to form according to the above analyses and the experimental results previously noted. In order to form a continuous SiC coating on HTGR graphite, we tried to reduce the pores size on the surface of the graphite through repeatedly infiltrating and pyrolyzing of the phenolic resin or by coating the substrate with nitrocellulose. The latter was found to be more effective. The nitrocellulose was first dissolved in acetone, and then it was coated on the substrate. A uniform and smooth coating formed on the surface of the substrate after pyrolysis of the cellulose. The surface pore size could be reduced to the range of 0.3–1.1 mm with only one coating-pyrolysis cycle. A continuous SiC coating was successfully formed after reducing the surface pore size by nitrocellulose pre-coating, supporting the coating formation mechanism noted above.
3.2. Improving oxidation resistance The oxidation resistance of SiC coating directly obtained from the reaction-formed process is insufficient because the coating normally contains many pinholes, pores and sometimes even cracks [11,15,16]. As seen in Fig. 3, SiC coatings obtained from the reaction-formed process are not fully dense. Many pinholes existed in the coatings. Pores up to 50 mm were frequently observed in the coating. Analyses showed that large pores mainly came from the substrate, which also had large pores on the surface. In this research, three measures were employed to improve the oxidation resistance of the coating. Namely, a primary coating was adopted to fill large voids on the surface of the substrate, free silicon was used to fill the pinholes and
Fig. 2. Pore size distribution of two kinds of graphite. The figure shows (A) cumulative pore volume vs. pore diameter and (B) incremental volume vs. pore diameter.
gradient SiC / C layer was applied to relax the thermal stress, thus reducing the possibility of crack generation.
3.2.1. Primary coating A primary coating which mainly consisted of SiC and graphite was coated on the substrate to fill the large voids (see Fig. 1). The porosity of the primary coating, which was designed to fit the volume expansion when the graphite was converted into SiC in the primary coating, can be controlled by adjusting the concentration of the slip. Experiments showed that the primary coating was quite effective in eliminating the large pores. Large pores were
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pinholes because free silicon was found to be quite effective in filling the pinholes. A dense SiSiC coating can be obtained directly, which would greatly simplify the whole process. Although free silicon may degrade the high temperature strength of the coating, in some applications such as the oxidation protection of fuel elements of HTGR under extremely accidental conditions (like large amount of water or air ingress into the reactor core) [30], the good oxidation resistance and low cost of the coating are of importance, whereas high temperature fracture strength is not critical. Sintering temperature had a big influence on the silicon content in the coating. As seen in Fig. 4, a weak peak of graphite was detected in the coating of 15008C, and no Si peaks were found below 16008C. The silicon peaks appeared from around 17008C, and the intensity of the Si peaks increased with increasing sintering temperature. The silicon content was quantitatively measured and is shown in Fig. 5. It shows that the free silicon content is |1.8% (volume percent) at 15008C and 16008C. As noted above, no free silicon was detected by SEM / EPMA and XRD over this temperature range; therefore, possibly this manifested the amount of closed pores in the coating. The amount of free Si increases with increasing temperature, consistent with the XRD results. The coating at 19508C contains about 5.5% silicon, and it is totally dense as shown in Fig. 6. Investigations showed that the coating could only be fully dense when it was sintered above
Fig. 3. Surface morphology of SiC coating obtained from reaction-formed process. The coating was made (a) at 16008C for 2 h and (b) at 18008C for 2 h.
rarely observed on the SiC coating when the primary coating was adopted.
3.2.2. Eliminating the pinholes As previously noted, the SiC coating obtained from the reaction-formed process contained many pinholes. Excellent oxidation resistance cannot be achieved by the coating, because oxygen can attack the substrate through the pinholes. The defects were filled with the primary SiO 2 coating and an enhanced coating by Shuford et al. [14]. Yamamot et al. [15] improved the oxidation resistance of the coating with an enhanced zircon coating obtained from a sol-gel process. CVD process was also used to improve the oxidation resistance of the coating [11]. Although good oxidation resistance was obtained with these enhanced coatings, the processes for the enhanced coatings were usually more complicated and required longer time, compared with the original reaction-formed process. In the present study, silicon was selected as the filler to fill these
Fig. 4. X-ray diffractograms of different SiC coatings. The figure shows that the intensity of Si peaks increases with increasing sintering temperature.
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Fig. 5. The influence of sintering temperature on the amount of free silicon in the coatings. The coatings were sintered for 2 h at different temperature.
19008C. It is easier for molten silicon to enter micropores, due to the increase in wettability and decrease in viscosity of molten silicon with increasing temperature [21], which is probably the reason for the increasing free silicon in the coating with temperature. The wettability of molten silicon on SiC can be enhanced by the addition of additives to the melt, which makes it possible to produce a dense SiSiC coating at a lower temperature. The addition of boron was reported to improve the wettability between molten silicon and reaction-formed SiC [31], but boron is a thermal
Fig. 6. Surface morphology of SiSiC coating sintered at 19508C for 2 h. The coating is totally dense, no pores and pinholes were observed in the coating.
Fig. 7. The influence of MgO addition on the amount of free silicon in the coatings. Coatings were sintered at 17008C for 2 h.
neutron poison, thus it was not used in this research. Experiments showed that the addition of MgO also had an effect similar to that of boron addition. Fig. 7 shows the free silicon content vs. the amount of MgO addition at 17008C. This shows that the amount of free silicon increases with increasing MgO addition. The coating contained |5.8% free silicon when it was sintered for 2 h at 17008C with the addition of 2.5 wt.% MgO. Without MgO addition, it needs to be sintered above 19008C to obtain the same amount of free silicon. Therefore, a dense SiSiC coating with excellent oxidation resistance can be obtained at a relatively low temperature with the addition of MgO.
3.2.3. Gradient SiC /C layer A SiC coating without a gradient SiC / C layer tends to crack under thermal cycling conditions, due to the thermal stress caused by the difference in thermal expansion between the SiC coating and the substrate [2,11]. Therefore, a dense SiC coating with a gradient SiC / C layer is essential to achieve good oxidation resistance, together with sufficient thermal shock resistance [6,15]. Although it is difficult to precisely control the SiC concentration in the gradient layer by the reaction-formed process, a dense SiC coating with a proper gradient SiC / C layer can be obtained through carefully controlling the sintering conditions [15]. Experiments showed that low temperature sintering (temperature around the melting point of silicon) was helpful to form the gradient SiC / C layer. For example, if the temperature was increased directly from room temperature to 17008C (no holding stage at around 14108C), as shown in Fig. 8a, there was a clear interface between the SiC
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3.3. Oxidation behavior 3.3.1. Low-temperature tests Low-temperature oxidation tests were performed on asreceived graphite, specimens with a porous SiC coating (coating with pinholes and pores) and specimens with a dense SiSiC coating from 8008C to 10008C. Fig. 9 shows the experimental results at 8008C. The as-received graphite loses |25% of its initial weight within 0.5 h. A remarkable improvement in oxidation resistance was observed on the sample with a porous SiC coating. The weight-loss was less than 10% during the same period. No weight loss was detected for dense SiSiC coating, indicating good oxidation resistance of the coating. It is known [32–34] that the oxidation of SiC at high oxygen partial pressure takes place according to reaction (3). Thermodynamic calculations [35] revealed that the oxidation behavior of the SiSiC coating was similar to that of the SiC coating, in which Si is simultaneously oxidized with SiC, as indicated by reactions (3) and (4). The transition behavior of SiSiC coating from passive oxidation to active oxidation was also shown to be similar to that of the SiC coating. SiC(S) 1 ]32 O 2 (g) → SiO 2 (S) 1 CO(g)
(3)
Si(S) 1 O 2 (g) → SiO 2 (S)
(4)
The oxidation rates were calculated from the weight loss data and are shown in Fig. 10. The oxidation rate of graphite increases with time, primarily due to the increase in surface area caused by the oxidation of graphite. The oxidation rate of the specimen with porous SiC coating changes little with time, as shown in Fig. 10. In this case, Fig. 8. Characteristic X-ray images of Si in the cross section of samples prepared by (a) directly sintering at 17008C for 2 h and (b) sintering at 14108C for 30 min, followed by sintering at 17008C for 1?5 h.
coating and the substrate. Normally, it took only 10 min to reach 17008C from 14008C; therefore, the SiC coating was mainly formed at high temperature. No gradient SiC / C layer formed in this case, possibly because the paths of infiltration were quickly blocked due to the rapid reaction of molten silicon with the substrate at 17008C. If, on the other hand, the temperature was held for 30 min at 14108C prior to 17008C, there is a gradual change in SiC concentration in the coating as shown in Fig. 8b, and no clear interface exists between the SiC coating and the substrate. The gradient SiC / C layer was believed to be formed mainly during the holding period at low temperature, where the reaction rate was relatively low. Molten silicon had sufficient time to penetrate into the interior of the substrate before the infiltration path was totally blocked, thus it is helpful for the formation of the gradient SiC / C layer.
Fig. 9. Oxidation-induced weight-loss vs. oxidation time.
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The small weight gain was caused by the formation of SiO 2 , due to the oxidation of SiC on the surface [34]. Low-temperature oxidation tests indicated that the SiSiC coating was fully dense, and no defects (cracks, pinholes or pores) existed in the coating. Otherwise, weight loss would be manifested, because these defects cannot be healed during oxidation at such a low temperature.
the oxidation process is limited by the diffusion of oxygen through the porous SiC coating. Because the porosity of the coating does not change much during oxidation, the overall reaction rate does not change much either. The oxidation rate of the SiSiC coating is virtually zero. Long time oxidation tests were performed on the samples with gradient SiSiC coating at 10008C. Fig. 11 shows the weight loss vs. oxidation time of a sample. No weight loss was observed after oxidizing for nearly 300 h.
3.3.2. High-temperature oxidation tests Thermal cyclic tests were carried out on a dense SiSiC coating with a gradient SiC / C layer at 12008C, 14008C and 16008C. As shown in Fig. 12, there was a small weight gain and no weight loss was found after oxidation for 1000 h at 12008C, together with 100 thermal cycles. The results of thermal cyclic tests at 14008C and 16008C are illustrated in Fig. 13. The samples were totally intact after oxidizing for 200 h at 14008C (20 thermal cycles) and 20 h at 16008C (four thermal cycles). The mass gain was due to the formation of SiO 2 . After thermal cyclic tests, no cracking or chipping off was found on the surface of the samples, indicating excellent oxidation resistance and thermal shock resistance of the coating. The good oxidation resistance is attributed to complete denseness of the coating. The high thermal shock resistance seems to be due to the gradient SiC / C layer which mitigates the thermal stress caused by the difference in thermal expansion between the SiC coating and the substrate [1,6]. The XRD pattern of the surface of the sample oxidized for 200 h at 14008C is shown in Fig. 14. It shows that the silica formed during oxidation is crystobalite. Strong peaks of b-SiC still exist in the pattern, suggesting that the silica film formed is very thin.
Fig. 11. Mass change plotted as a function of oxidation time. Negative weight-loss means mass gain. The specimen was isothermally oxidized at 10008C.
Fig. 12. Mass change plotted as a function of oxidation time. Negative weight-loss means mass gain. The specimen was tested at 12008C for 1000 h together with 100 thermal cycles.
Fig. 10. Oxidation rate vs. oxidation time at 8008C. The oxidation rate was calculated from the weight loss data in Fig. 9.
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reaction occurs not only on the surface but also inside the substrate, which results in a big volume expansion. Cracks and spallation are thus inevitable. Theoretical analyses showed that the infiltration rate could be controlled through adjusting the pore size on the surface of the substrate, consistent with the experimental results. 3. Oxidation tests showed that SiSiC coatings were fully dense. Samples were totally intact after oxidizing over 1000 h with 100 thermal cycles at 12008C and 200 h with 20 thermal cycles at 14008C, demonstrating excellent oxidation resistance and thermal shock resistance. This high thermal shock resistance is attributed to the gradient SiC / C layer which could mitigate the thermal stress between the SiC coating and the substrate.
References Fig. 13. Mass change plotted as a function of oxidation time. Negative weight-loss means mass gain. The specimens were tested at 14008C and 16008C under thermal cyclic conditions.
4. Conclusions
1. A dense functional gradient SiSiC coating could be obtained directly from a novel process, in which large pores in the coating were eliminated by a primary coating and the pinholes were filled with free silicon. The amount of free silicon in the coating can be adjusted by changing the reaction temperature and adjusting the wettability of molten silicon with additives. 2. The investigations revealed that the competition between the infiltration and the reaction determined whether a continuous SiC coating was formed. If the infiltration rate is equal to or less than the reaction rate, the infiltration and the reaction are limited only to the surface of the substrate. A continuous SiC coating can be formed by the reaction. However, if the infiltration speed is much higher than the reaction speed, the
Fig. 14. XRD pattern of the surface of the sample oxidized for 200 h at 14008C.
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Oxidation resistant SiC coating for graphite materials Qingshan Zhu*, Xueliang Qiu, Changwen Ma Institute of Nuclear Energy Technology, Tsinghua University, Beijing 100080, People’ s Republic of China Received 25 November 1998; accepted 8 January 1999
Abstract A dense functional gradient SiSiC coating with good oxidation resistance when applied to a graphite substrate has been developed using a novel process, in which large pores in the coating were eliminated by application of a primary coating, and pinholes were filled with free silicon. The mechanism for the formation of a continuous SiC coating on graphite was investigated experimentally and theoretically. Oxidation tests were conducted from 8008C to 16008C for up to 1000 h. Samples of graphite with the SiSiC coating were found to be totally intact after oxidizing for 1000 h together with 100 thermal cycles at 12008C and 200 h with 20 thermal cycles at 14008C, indicating excellent oxidation resistance and thermal shock resistance of the coating. The good thermal shock resistance is attributed to the compositional gradient in the coating. 1999 Elsevier Science Ltd. All rights reserved. Keywords: A. Graphite; B. Coating; Oxidation
1. Introduction Carbon and graphite are attractive materials for high temperature applications due to their high strength, high modulus, excellent thermal shock resistance and light weight. They are widely used as engineering materials, such as heaters, electrical contacts, high-temperature heat exchangers, rocket nozzles and leading edges of aircraft wings, etc. [1,2]. However, the use of carbon materials has been greatly restricted due to the poor oxidation resistance at high temperature in an oxidizing atmosphere. Achieving good oxidation resistance is crucial to utilize their full potential as high-temperature materials. Oxidation protection for carbon materials has been extensively studied in the past 60 years [3–5]. Ceramic coatings are commonly employed to protect carbon materials from oxidation [6–8]. Although several coating systems have been developed, SiC is considered to be the best coating material due to its good mechanical properties, low density, good physical–chemical compatibility with carbon and excellent oxidation resistance below 18008C [9]. The SiC coating can be formed by several methods [10–13]. Reaction-formed process, in which molten silicon reacts with the substrate on the surface to form an SiC coating, is one of the effective ways to form an SiC coating [14]. However, the SiC coating obtained from reaction-formed *Corresponding author.
process often contains many defects (pores, pinholes or cracks), and the oxidation resistance of the coating is not sufficient at high temperature, because oxygen can attack the substrate through these defects. Additional processes were usually needed to fill these defects. These processes are complicated and required longer time than the original reaction-formed process [15,16]. Another problem is that the coating formation behavior can be quite different when different carbon materials are used; sometimes it is difficult to form an SiC coating by this process [17]. Therefore, the coating formation mechanism still needs to be clarified. In the present work, the mechanism for the formation of the SiC coating was systematically studied. A new process, which was aimed at making a dense SiC coating directly, was developed and optimized. A dense SiSiC coating with a gradient SiC / C layer was formed on two kinds of graphite. The oxidation behavior of the coatings was tested from 8008C to 16008C for up to 1000 h.
2. Experimental
2.1. Materials Two kinds of graphite were used in this research. One was the matrix graphite of spherical fuel elements of a high temperature gas cooled reactor (HTGR), which is
0008-6223 / 99 / $ – see front matter 1999 Elsevier Science Ltd. All rights reserved. PII: S0008-6223( 99 )00010-X
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under construction at the authors’ institute. The matrix graphite is made of 64 wt.% natural flake graphite, 16 wt.% petroleum coke graphite and 20 wt.% resin binder [18]. This graphite will be termed ‘HTGR graphite’ hereafter. The other was a type of nuclear-grade graphite provided by Shanghai Carbon Corporation (Shanghai, P.R. China). It will be termed ‘Shanghai graphite’ in the rest of this paper. The characteristics of the graphite materials are summarized in Table 1. The graphite was cut into small specimens with a size of 2031538 mm. The surface of the specimens was polished with 5.0 mm diamond paste, followed by ultrasonically washing in acetone for 30 min. Other materials used in the process were Si powder (circa 1 mm), SiC powder (a, circa 0.1 mm) and nucleargrade graphite powder (circa 1 mm). Several reagents were also employed in the slip process. These reagents were (1) the acrylamide [CH 2 =CHCONH 2 ], the dispersant and monomer; (2) N,N9-methylene bisacrylamide [(CH 2 =CHCONH) 2 CH 2 ], the crosslinking agent; (3) ammonium persulfate [(NH 4 ) 2 S 2 O 8 ], the free-radical initiator for polymerization; and (4) tetramethylammonium hydroxide [(CH 3 ) 4 NOH], the agent for adjusting the pH value of the slip.
2.2. Coating preparation procedure A new process for preparing an oxidation resistant coating for graphite materials was developed from the conventional reaction-formed process [16] and the gelcasting process [19]. The preparation procedure of the process is shown in Fig. 1. The process consisted of three steps. First, a primary SiC coating was coated on the surface of a substrate, then the substrate with the primary coating was encased in a slip of silicon to form a packed gel, and finally, the packed body was sintered to produce the SiC coating. To produce a primary coating on a substrate, a-SiC powder, graphite powder and water were mixed with a specified amount of dispersant and crosslinking agent to form a slurry. The pH value of the slurry was adjusted to 10–12. The slip was milled for a sufficient time before the addition of the initiator. The slip was then sprayed on the surface of the substrate. The monomer in the slip polymerized to form a primary SiC coating, which could attach strongly to the substrate after drying. A packing slip, which consisted of Si powder, a-SiC
Table 1 Characteristics of graphite materials Graphite
Shanghai HTGR
r kg / m 3
1810 1730
Ash content 310 6
400 130–190
Thermal expansion coefficient 310 6 / K '
i
5.50 2.77
4.27 2.54
Fig. 1. Coating preparation procedure.
powder and water, was mixed and milled with a specified amount of dispersant and crosslinking agent. The substrate with primary coating was then dipped into the packing slip, and in the meantime, the initiator was added to the packing slip to initiate the polymerization reaction. A packed body, which contained the substrate, formed after the polymerization of the monomer. The speed of polymerization could be controlled by adjusting the temperature and the concentration of the initiator and crosslinking agent. The packed body was dried in an oven between 508C to 2008C. No special care is needed in the drying process. The dried body has very high strength, for example, it can even be machined if desired [19]. Consequently, there are no handling problems with the dried body. The oxidation resistant SiC coating could be obtained by sintering the packing body to a temperature between 15008C and 19508C for 1 to 4 h.
2.3. Coating characterization The surface of the coatings was observed by optical microscopy and scanning electron microscopy (SEM). Coating quality was assessed through examining the surface for cracks, pores and pinholes. Elemental distribution of Si in the cross section of the coatings was detected by electron probe microanalysis (EPMA) to study the distribution of the SiC phase. The crystalline phases of the coatings were evaluated by X-ray diffractometer (XRD). The density of the SiSiC coating, which was used to calculate the amount of free silicon in the coating through Eq. (1) by assuming that the coating was composed only of Si and SiC, was measured by a titration method [20].
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rm 5 rSi x 1 rSiC (1 2 x)
(1)
where rm is the measured density of the SiSiC coating, rSi the density of Si, rSiC the density of SiC and x the volume percentage of Si in the SiSiC coating. Before measuring the density, the SiSiC coating was burned at 6008C for a sufficient time to remove the attached graphite. The burned coating was ultrasonically washed in acetone for a sufficient time to remove the ash, followed by drying in an oven. Diiodomethane (CH 2 I 2 ) and bromoform (CHBr 3 ) were used as the titrants. The density of diiodomethane and bromoform is 3.326310 3 kg / m 3 , 2.893310 3 kg / m 3 and 3.320310 3 kg / m 3 , 2.889310 3 kg / m 3 at 48C and 208C, respectively.
2.4. Oxidation tests Thermogravimetric tests were carried out under natural convection conditions in air between 8008C and 10008C. The sample was supported by a silica basket hung on a silica chain from a Mettler balance. The sensitivity of the balance is 0.1 mg. The mass change was continuously recorded by a PC computer. Thermal cyclic tests at high temperature were carried out in a MoSi 2 furnace. The furnace was first heated to a specified temperature (12008C, 14008C and 16008C), then a sample was placed into the furnace and held for a specified time (10 h for the tests at 12008C and 14008C, and 5 h for 16008C). After oxidation, the sample was taken out and cooled at room temperature. The surface temperature of the sample can fall down to 2008C within 5 min. After cooling to room temperature, the sample was weighed and placed in the furnace again.
3. Results and discussion
3.1. Coating formation mechanism The reaction mechanism between solid carbon and molten silicon has been studied by various authors [21– 23]. Although SiC was proved to be formed through a solution–reprecipitation mechanism on a micro-scale, different results can be generated in forming the SiC coating by the reaction. For example, both Minnear [24] and Yamamoto et al. [15] reported that a continuous SiC layer could be formed by this reaction, while Hase and Suzuki [17] found no continuous SiC layer between the carbon particles and silicon and suggested that the reaction product quickly spalled due to the volume misfit between SiC and the carbon [22,25]. In our experiments, the authors also found that a continuous SiC coating was easy to form on Shanghai graphite, while it was difficult to form a continuous SiC coating on HGTR graphite by this reaction. In the case of HTGR graphite, big cracks and severe spallation occurred. Sometimes the substrate was crushed
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into powder by the reaction. The volume will expand |234% [26] when the graphite is totally converted into SiC. This volume expansion is believed to be the reason for the cracks and spallation, but the volume misfit of Shanghai graphite should be larger than that of HTGR graphite, because the density of Shanghai graphite is greater than that of HTGR graphite. Various investigations have been conducted to find the direct reason for this. It was found that the competition between the infiltration process and the chemical reaction process determined whether a continuous SiC coating was formed. If the infiltration rate is smaller than that of the reaction rate, the whole process is controlled by the infiltration process. The infiltrant (molten silicon) will be consumed at the infiltration frontier. During infiltration, the reaction-induced volume expansion will reduce the paths of infiltration, which further decreases the infiltration rate. Finally, the paths of infiltration will be totally blocked by the volume expansion and a continuous SiC coating forms. After that, the growth of the SiC layer is controlled by the diffusion of Si and C through the newly formed SiC layer. In this case, the reaction is limited only to the surface of the substrate. Therefore, the total volume expansion is small. If, on the other hand, the infiltration rate is much higher than the chemical reaction rate, the infiltrant can be delivered rapidly ahead of the reaction front. A large amount of molten silicon will penetrate deeply into the substrate before the infiltration paths are narrowed by reaction-induced volume expansion. The substrate will be converted into SiC not only on the surface but also inside the substrate, which results in a big volume expansion. This big volume expansion will cause cracks or spallation of the substrate even if a layer of SiC has already formed on the surface of the substrate. Cracks may also be caused by the thermal stress resulting from rapid local heating due to the quick reaction. The cracks can serve as the new paths for infiltration. More infiltrant can penetrate into the substrate through the newly formed cracks, and a big volume expansion can be generated by the consequent reaction, which may cause new cracks again. Such cycles can continue until the substrate is completely converted into SiC. The chemical reaction rate is a function of temperature and the graphite. The infiltration speed v can be expressed by the following equation [27]: dl rgl n cosu n 5 ] 5 ]]] dt 4hl
(2)
where t is the infiltration time, l the length of infiltration at time t, r the equivalent radius of the pore, glv the surface tension, u the wetting angle and h the viscosity of the infiltrant. When the system and experimental conditions are fixed, the chemical reaction rate is fixed, and glv , u and h are also fixed. The rate of infiltration is proportional to the pore diameter as indicated by Eq. (2). Therefore, the
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competition of these two processes is mainly determined by the pore size of the substrate. This point was confirmed by the experimental results of Singh and Behrendt [28], which revealed that the speed of infiltration and reaction was almost the same in the case of a small pore size (0.2 mm), while in the case of a medium pore size (1.8 mm), the infiltration speed was much higher than the reaction rate. Our investigations showed that the difference between Shanghai graphite and HTGR graphite in coating formation mainly lay in the pore structure, because the chemical kinetics of the two kinds of graphite with molten silicon was almost the same, as confirmed by differential thermal analyses (DTA) [29]. The pore size distribution of the two kinds of graphite is shown in Fig. 2. The pore size of Shanghai graphite is mainly over the range of 0.3–0.9 mm, which is very close to the small pore size of 0.2 mm examined by Singh and Behrendt [28]. In this case, the speed of the infiltration and the reaction must be in the same order, and a continuous SiC coating is expected to form. As seen in Fig. 2, the pore size of HTGR graphite shows a bimodal distribution, and the pores are mainly in the range of 2–4 mm, which is larger than the medium pore size indicated by Singh and Behrendt [28]. In this case, the infiltration speed must be much higher than the reaction rate. A continuous SiC coating is difficult to form according to the above analyses and the experimental results previously noted. In order to form a continuous SiC coating on HTGR graphite, we tried to reduce the pores size on the surface of the graphite through repeatedly infiltrating and pyrolyzing of the phenolic resin or by coating the substrate with nitrocellulose. The latter was found to be more effective. The nitrocellulose was first dissolved in acetone, and then it was coated on the substrate. A uniform and smooth coating formed on the surface of the substrate after pyrolysis of the cellulose. The surface pore size could be reduced to the range of 0.3–1.1 mm with only one coating-pyrolysis cycle. A continuous SiC coating was successfully formed after reducing the surface pore size by nitrocellulose pre-coating, supporting the coating formation mechanism noted above.
3.2. Improving oxidation resistance The oxidation resistance of SiC coating directly obtained from the reaction-formed process is insufficient because the coating normally contains many pinholes, pores and sometimes even cracks [11,15,16]. As seen in Fig. 3, SiC coatings obtained from the reaction-formed process are not fully dense. Many pinholes existed in the coatings. Pores up to 50 mm were frequently observed in the coating. Analyses showed that large pores mainly came from the substrate, which also had large pores on the surface. In this research, three measures were employed to improve the oxidation resistance of the coating. Namely, a primary coating was adopted to fill large voids on the surface of the substrate, free silicon was used to fill the pinholes and
Fig. 2. Pore size distribution of two kinds of graphite. The figure shows (A) cumulative pore volume vs. pore diameter and (B) incremental volume vs. pore diameter.
gradient SiC / C layer was applied to relax the thermal stress, thus reducing the possibility of crack generation.
3.2.1. Primary coating A primary coating which mainly consisted of SiC and graphite was coated on the substrate to fill the large voids (see Fig. 1). The porosity of the primary coating, which was designed to fit the volume expansion when the graphite was converted into SiC in the primary coating, can be controlled by adjusting the concentration of the slip. Experiments showed that the primary coating was quite effective in eliminating the large pores. Large pores were
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pinholes because free silicon was found to be quite effective in filling the pinholes. A dense SiSiC coating can be obtained directly, which would greatly simplify the whole process. Although free silicon may degrade the high temperature strength of the coating, in some applications such as the oxidation protection of fuel elements of HTGR under extremely accidental conditions (like large amount of water or air ingress into the reactor core) [30], the good oxidation resistance and low cost of the coating are of importance, whereas high temperature fracture strength is not critical. Sintering temperature had a big influence on the silicon content in the coating. As seen in Fig. 4, a weak peak of graphite was detected in the coating of 15008C, and no Si peaks were found below 16008C. The silicon peaks appeared from around 17008C, and the intensity of the Si peaks increased with increasing sintering temperature. The silicon content was quantitatively measured and is shown in Fig. 5. It shows that the free silicon content is |1.8% (volume percent) at 15008C and 16008C. As noted above, no free silicon was detected by SEM / EPMA and XRD over this temperature range; therefore, possibly this manifested the amount of closed pores in the coating. The amount of free Si increases with increasing temperature, consistent with the XRD results. The coating at 19508C contains about 5.5% silicon, and it is totally dense as shown in Fig. 6. Investigations showed that the coating could only be fully dense when it was sintered above
Fig. 3. Surface morphology of SiC coating obtained from reaction-formed process. The coating was made (a) at 16008C for 2 h and (b) at 18008C for 2 h.
rarely observed on the SiC coating when the primary coating was adopted.
3.2.2. Eliminating the pinholes As previously noted, the SiC coating obtained from the reaction-formed process contained many pinholes. Excellent oxidation resistance cannot be achieved by the coating, because oxygen can attack the substrate through the pinholes. The defects were filled with the primary SiO 2 coating and an enhanced coating by Shuford et al. [14]. Yamamot et al. [15] improved the oxidation resistance of the coating with an enhanced zircon coating obtained from a sol-gel process. CVD process was also used to improve the oxidation resistance of the coating [11]. Although good oxidation resistance was obtained with these enhanced coatings, the processes for the enhanced coatings were usually more complicated and required longer time, compared with the original reaction-formed process. In the present study, silicon was selected as the filler to fill these
Fig. 4. X-ray diffractograms of different SiC coatings. The figure shows that the intensity of Si peaks increases with increasing sintering temperature.
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Fig. 5. The influence of sintering temperature on the amount of free silicon in the coatings. The coatings were sintered for 2 h at different temperature.
19008C. It is easier for molten silicon to enter micropores, due to the increase in wettability and decrease in viscosity of molten silicon with increasing temperature [21], which is probably the reason for the increasing free silicon in the coating with temperature. The wettability of molten silicon on SiC can be enhanced by the addition of additives to the melt, which makes it possible to produce a dense SiSiC coating at a lower temperature. The addition of boron was reported to improve the wettability between molten silicon and reaction-formed SiC [31], but boron is a thermal
Fig. 6. Surface morphology of SiSiC coating sintered at 19508C for 2 h. The coating is totally dense, no pores and pinholes were observed in the coating.
Fig. 7. The influence of MgO addition on the amount of free silicon in the coatings. Coatings were sintered at 17008C for 2 h.
neutron poison, thus it was not used in this research. Experiments showed that the addition of MgO also had an effect similar to that of boron addition. Fig. 7 shows the free silicon content vs. the amount of MgO addition at 17008C. This shows that the amount of free silicon increases with increasing MgO addition. The coating contained |5.8% free silicon when it was sintered for 2 h at 17008C with the addition of 2.5 wt.% MgO. Without MgO addition, it needs to be sintered above 19008C to obtain the same amount of free silicon. Therefore, a dense SiSiC coating with excellent oxidation resistance can be obtained at a relatively low temperature with the addition of MgO.
3.2.3. Gradient SiC /C layer A SiC coating without a gradient SiC / C layer tends to crack under thermal cycling conditions, due to the thermal stress caused by the difference in thermal expansion between the SiC coating and the substrate [2,11]. Therefore, a dense SiC coating with a gradient SiC / C layer is essential to achieve good oxidation resistance, together with sufficient thermal shock resistance [6,15]. Although it is difficult to precisely control the SiC concentration in the gradient layer by the reaction-formed process, a dense SiC coating with a proper gradient SiC / C layer can be obtained through carefully controlling the sintering conditions [15]. Experiments showed that low temperature sintering (temperature around the melting point of silicon) was helpful to form the gradient SiC / C layer. For example, if the temperature was increased directly from room temperature to 17008C (no holding stage at around 14108C), as shown in Fig. 8a, there was a clear interface between the SiC
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3.3. Oxidation behavior 3.3.1. Low-temperature tests Low-temperature oxidation tests were performed on asreceived graphite, specimens with a porous SiC coating (coating with pinholes and pores) and specimens with a dense SiSiC coating from 8008C to 10008C. Fig. 9 shows the experimental results at 8008C. The as-received graphite loses |25% of its initial weight within 0.5 h. A remarkable improvement in oxidation resistance was observed on the sample with a porous SiC coating. The weight-loss was less than 10% during the same period. No weight loss was detected for dense SiSiC coating, indicating good oxidation resistance of the coating. It is known [32–34] that the oxidation of SiC at high oxygen partial pressure takes place according to reaction (3). Thermodynamic calculations [35] revealed that the oxidation behavior of the SiSiC coating was similar to that of the SiC coating, in which Si is simultaneously oxidized with SiC, as indicated by reactions (3) and (4). The transition behavior of SiSiC coating from passive oxidation to active oxidation was also shown to be similar to that of the SiC coating. SiC(S) 1 ]32 O 2 (g) → SiO 2 (S) 1 CO(g)
(3)
Si(S) 1 O 2 (g) → SiO 2 (S)
(4)
The oxidation rates were calculated from the weight loss data and are shown in Fig. 10. The oxidation rate of graphite increases with time, primarily due to the increase in surface area caused by the oxidation of graphite. The oxidation rate of the specimen with porous SiC coating changes little with time, as shown in Fig. 10. In this case, Fig. 8. Characteristic X-ray images of Si in the cross section of samples prepared by (a) directly sintering at 17008C for 2 h and (b) sintering at 14108C for 30 min, followed by sintering at 17008C for 1?5 h.
coating and the substrate. Normally, it took only 10 min to reach 17008C from 14008C; therefore, the SiC coating was mainly formed at high temperature. No gradient SiC / C layer formed in this case, possibly because the paths of infiltration were quickly blocked due to the rapid reaction of molten silicon with the substrate at 17008C. If, on the other hand, the temperature was held for 30 min at 14108C prior to 17008C, there is a gradual change in SiC concentration in the coating as shown in Fig. 8b, and no clear interface exists between the SiC coating and the substrate. The gradient SiC / C layer was believed to be formed mainly during the holding period at low temperature, where the reaction rate was relatively low. Molten silicon had sufficient time to penetrate into the interior of the substrate before the infiltration path was totally blocked, thus it is helpful for the formation of the gradient SiC / C layer.
Fig. 9. Oxidation-induced weight-loss vs. oxidation time.
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The small weight gain was caused by the formation of SiO 2 , due to the oxidation of SiC on the surface [34]. Low-temperature oxidation tests indicated that the SiSiC coating was fully dense, and no defects (cracks, pinholes or pores) existed in the coating. Otherwise, weight loss would be manifested, because these defects cannot be healed during oxidation at such a low temperature.
the oxidation process is limited by the diffusion of oxygen through the porous SiC coating. Because the porosity of the coating does not change much during oxidation, the overall reaction rate does not change much either. The oxidation rate of the SiSiC coating is virtually zero. Long time oxidation tests were performed on the samples with gradient SiSiC coating at 10008C. Fig. 11 shows the weight loss vs. oxidation time of a sample. No weight loss was observed after oxidizing for nearly 300 h.
3.3.2. High-temperature oxidation tests Thermal cyclic tests were carried out on a dense SiSiC coating with a gradient SiC / C layer at 12008C, 14008C and 16008C. As shown in Fig. 12, there was a small weight gain and no weight loss was found after oxidation for 1000 h at 12008C, together with 100 thermal cycles. The results of thermal cyclic tests at 14008C and 16008C are illustrated in Fig. 13. The samples were totally intact after oxidizing for 200 h at 14008C (20 thermal cycles) and 20 h at 16008C (four thermal cycles). The mass gain was due to the formation of SiO 2 . After thermal cyclic tests, no cracking or chipping off was found on the surface of the samples, indicating excellent oxidation resistance and thermal shock resistance of the coating. The good oxidation resistance is attributed to complete denseness of the coating. The high thermal shock resistance seems to be due to the gradient SiC / C layer which mitigates the thermal stress caused by the difference in thermal expansion between the SiC coating and the substrate [1,6]. The XRD pattern of the surface of the sample oxidized for 200 h at 14008C is shown in Fig. 14. It shows that the silica formed during oxidation is crystobalite. Strong peaks of b-SiC still exist in the pattern, suggesting that the silica film formed is very thin.
Fig. 11. Mass change plotted as a function of oxidation time. Negative weight-loss means mass gain. The specimen was isothermally oxidized at 10008C.
Fig. 12. Mass change plotted as a function of oxidation time. Negative weight-loss means mass gain. The specimen was tested at 12008C for 1000 h together with 100 thermal cycles.
Fig. 10. Oxidation rate vs. oxidation time at 8008C. The oxidation rate was calculated from the weight loss data in Fig. 9.
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reaction occurs not only on the surface but also inside the substrate, which results in a big volume expansion. Cracks and spallation are thus inevitable. Theoretical analyses showed that the infiltration rate could be controlled through adjusting the pore size on the surface of the substrate, consistent with the experimental results. 3. Oxidation tests showed that SiSiC coatings were fully dense. Samples were totally intact after oxidizing over 1000 h with 100 thermal cycles at 12008C and 200 h with 20 thermal cycles at 14008C, demonstrating excellent oxidation resistance and thermal shock resistance. This high thermal shock resistance is attributed to the gradient SiC / C layer which could mitigate the thermal stress between the SiC coating and the substrate.
References Fig. 13. Mass change plotted as a function of oxidation time. Negative weight-loss means mass gain. The specimens were tested at 14008C and 16008C under thermal cyclic conditions.
4. Conclusions
1. A dense functional gradient SiSiC coating could be obtained directly from a novel process, in which large pores in the coating were eliminated by a primary coating and the pinholes were filled with free silicon. The amount of free silicon in the coating can be adjusted by changing the reaction temperature and adjusting the wettability of molten silicon with additives. 2. The investigations revealed that the competition between the infiltration and the reaction determined whether a continuous SiC coating was formed. If the infiltration rate is equal to or less than the reaction rate, the infiltration and the reaction are limited only to the surface of the substrate. A continuous SiC coating can be formed by the reaction. However, if the infiltration speed is much higher than the reaction speed, the
Fig. 14. XRD pattern of the surface of the sample oxidized for 200 h at 14008C.
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