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Ablation behavior of ZrB2–SiC–ZrO2 ceramic composites by means of the oxyacetylene torch
Corrosion Science 51 (2009) 2071–2079
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Ablation behavior of ZrB2–SiC–ZrO2 ceramic composites by means of the oxyacetylene torch Shanbao Zhou a, Weijie Li a,*, Ping Hu a, Changqing Hong a, Ling Weng b a b
Center for Composite Materials, Harbin Institute of Technology, Harbin 150001, PR China School of Material Science and Engineering, Harbin University of Science and Technology, Harbin 150001, PR China
a r t i c l e
i n f o
Article history: Received 9 March 2009 Accepted 27 May 2009 Available online 3 June 2009 Keywords: A. Ceramic matrix composites B. SEM B. Weight loss C. Oxidation
a b s t r a c t Ablation behavior of ZrB2–SiC–ZrO2 ceramics with two ZrO2 contents was investigated using oxyacetylene torch. Thermogravimetric analysis demonstrated that ceramic with 10 vol% ZrO2 showed initial weight change at higher temperature than the one with 20 vol% ZrO2. After same ablation condition, lighter oxidized microstructure and lower weight loss and line gain were obtained from ceramic with 10 vol% ZrO2. Ablation mechanism revealed that excessive ZrO2 would supply much path to the inward transport of oxygen, which led to the dissatisfactory resistance to oxidation and ablation for the ceramic with 20 vol% ZrO2. Ó 2009 Elsevier Ltd. All rights reserved.
1. Introduction Transition metal borides including ZrB2, HfB2 and TiB2 are of great interest for their unique properties combined with ultra-high melting temperature (>3000 °C), high electrical and thermal conductivities, as well as excellent chemical inertness [1]. Such properties make them the potential candidates for leading edges on aircraft and re-entry vehicles, and for structural parts in high temperature environments [2,3]. Development of these ceramics has focused on densification [4,5], mechanical properties [5–7], thermal shock and oxidation behavior [8–13]. Monolithic ZrB2 was limited in the sinterability due to its covalent bonding, high melting temperature, and low self-diffusion coefficients of Zr and B. In this perspective, SiC was introduced to overcome this obstacle, and near theoretical density was achieved from ZrB2–SiC system. In addition, dissatisfactory toughness of ZrB2 ceramics was still a shortcoming to a wider range of use for most sophisticated applications [1]. Given that defect, phase transformation of ZrO2 was proposed to toughen ZrB2-based ceramics in the previous work [14,15]. As a critical enhancement, oxidation and ablation behavior was investigated widely by many researchers. In terms of the oxidation of ZrB2-based ceramics, ZrB2 is first oxidized to produce a scale composed of ZrO2 and B2O3 at 800 °C under the oxidizing conditions [16]. Due to the lower melting temperature (450 °C), liquid B2O3 can form a continuous layer that seals the surface and acts as an effective barrier to the inward transport of oxygen below * Corresponding author. Tel./fax: +86 451 86402382. E-mail address: [email protected] (W. Li). 0010-938X/$ - see front matter Ó 2009 Elsevier Ltd. All rights reserved. doi:10.1016/j.corsci.2009.05.035
1100 °C. At temperatures above 1100 °C, B2O3 begins to evaporate and such behavior becomes dominant when temperature is higher than 1400 °C due to the high vapor pressure of B2O3 [16–18]. However, the oxidation of ZrB2 can be improved by adding SiC to promote the formation of a silica-rich scale on the exposed surfaces with temperature above 1100 °C. That silica-rich layer is also refractory and then provides resistance to oxygen flow [9,16–18]. Moreover, such layer will resist oxidation effectively up to at least 1500 °C, because SiO2 is evidently less volatile than B2O3 at this temperature [17]. Once the temperature is above 2000 °C and even up to 2100 °C, the effective protection from SiC is obviously weakened due to the active oxidation of SiC to produce SiO, as well as the rapid evaporation of SiO2 via the high vapor pressure [9,19]. ZrO2 skeleton is then dominantly responsible for the oxidation resistance of ZrB2-based ceramics at this stage [9,19]. Under the ablation environment, mechanical loads are inflicted on the ceramics in addition to the oxidation. Much progress has been also carried out in ZrB2-based and ZrB2-conatining ceramics by means of the oxyacetylene torch and arc-heated wind tunnel tests [20–22]. Because of the mechanical loads-induced scouring, weight of the ceramics showed great variation compared with the sole oxidation process [20,22]. Analogously, microstructure evolution was rather complicated for ceramics after ablation, which provided typical layered microstructure along the cross-section [23], as well as the trace of mechanical scouring in ceramics [20–22]. Based on the reported work about ZrO2-toughened ZrB2 ceramics, the present paper was newly proposed to evaluate the ablation behavior of ZrB2–SiC–ZrO2 ceramics using oxyacetylene torch. Effect of ZrO2 content on the ablation behavior was determined in
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2. Experimental
where m1 was the weight (mg) of ceramic specimen before ablation, m2 was the one after ablation, A was the surface area (cm2) of the specimen exposed for ablation. The specific line gain (mm/s) was calculated by measuring the line before and after ablation with the following equation:
2.1. Starting powders and preparation
Dl ¼
Commercially available raw materials were used to prepare the ZrB2–SiC–ZrO2 composite in this work. The ZrB2 powder with a mean size of 2 lm was supplied from the Northwest Institute for Non-Ferrous Metal Research, China. The SiC particles (1 lm, Weifang Kaihua Micro-Powder Co., Ltd., Shandong, China) were predominantly a-SiC. The ZrO2 (1 lm, Fanmeiya Powders Co., Ltd., Jiangxi, China) used here was 3 mol% Y2O3 partially stabilized zirconia prepared by the co-precipitation method [24]. Two fractions of ZrO2, i.e., 10 vol% and 20 vol%, were considered to add to ZrB2–10 vol%SiC ceramics, and the two ceramics were marked as BS10Z and BS20Z, respectively. According to the detailed composition, powder mixtures were milled for 8 h in a polyethylene bottle with ZrO2 balls and ethanol as the grinding media. After mixing, the slurry was dried in a rotary evaporator and screened. Milled powders were hot-pressed at 1850 °C for 60 min under a uniaxial load of 30 MPa in Ar atmosphere.
where l1 and l2 were the height (mm) of the specimen before and after ablation, t was the ablation time (s).
terms of the ablation and oxidation mechanism with investigating the microstructure as well as the variation of weight and line after ablation.
2.2. Ablation tests Ablation behavior of the present ZrB2–SiC–ZrO2 ceramics was evaluated by means of the oxyacetylene torch. Table 1 lists the specific experimental condition and the measured surface temperature. During the ablation tests, a specimen with 15 mm in diameter and 10 mm in height was exposed to the flame. Ablation time was chosen as 30 s, 60 s, 120 s, 180 s, 240 s, 360 s and 600 s. The distance between the nozzle tip of the oxyacetylene gun and the surface of ceramic specimen was 10 mm. In order to ensure the ablation in just one direction, the specimen was held in a concave graphite anvil and only one face was exposed to the ablation environment. Flow rates of the oxygen and acetylene were 1.02 L/s and 0.36 L/s, respectively. The inner diameter of the nozzle tip of the ablation gun was 2 mm. Surface temperature of the ceramic samples was determined by means of an optical pyrometer and recorded every 10 s. 2.3. Characterization Crystalline phases of the specimen after ablation were identified by X-ray diffraction (XRD, Rigaku, Dmax-rb). The microstructure after ablation was analyzed by scanning electron microscopy/energy dispersive spectroscopy (SEM/EDS, Quanta 200). Thermogravimetric analysis (TGA) was conducted on a Netzsch STA 449C using a heating rate of 10 °C/min from 20 °C to 1400 °C. The specific weight loss (mg/cm2) was calculated by measuring the weight before and after ablation with the following equation:
Dm ¼
m1 m2 A
l2 l1 t
ð2Þ
3. Results and discussion 3.1. TGA TGA experiments were conducted firstly for the two ZrB2–SiC– ZrO2 ceramics from 20 °C (room temperature) to 1400 °C to assist exploring the ablation mechanism. As plotted in Fig. 1, the weight did not change with the increased testing temperature for both ceramics below 700 °C. The first time for weight change of ceramic BS10Z occurred at 795 °C, and BS20Z at 735 °C. Such weight change indicated the presence of the initial oxidation of ZrB2 in ZrB2–SiC–ZrO2 ceramics according to Reaction (3). As the temperature elevated, weight change increased gradually for both ceramics. Previous studies reported that B2O3 evaporated rapidly (Reaction (4)) at temperature above 1100 °C [16–19]. At 1150 °C in Fig. 1, an evident inflection point appeared for the weight change curve of ceramic BS20Z, and similar inflection point for BS10Z was at 1220 °C. Such point represented that SiC in ZrB2–SiC–ZrO2 ceramics began to be oxidized rapidly according to Reaction (5), which was basically consistent with the reported progress [16]. After that, the weight for both ceramics both provided monotonous increment with extended temperature.
ZrB2 (s) + 5/2O2 (g) = ZrO2 (s) + B2 O3 (l)
ð3Þ
B2 O3 (l) = B2 O3 (g)
ð4Þ
SiC(s) + 3/2O2 (g) = SiO2 (l) + CO(g)
ð5Þ
3.2. Macroscopical morphology Fig. 2 plots the temperature curves on the surfaces of the two ZrB2–SiC–ZrO2 ceramics varied with the ablation time increased up to 600 s. Sharp increment in the surface temperature could be
ð1Þ
Table 1 Experimental condition for oxyacetylene ablation and the measured surface temperature. O2 gas pressure (kPa) O2 gas flux (L/s) C2H2 gas pressure (kPa) C2H2 gas flux (L/s) Diameter of nozzle (mm) Distance from sample surface to nozzle (mm) Measured surface temperature (°c)
0.45 1.02 0.1 0.36 2 10 2000 ± 1
Fig. 1. TGA of BS10Z and BS20Z ceramics in air.
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Fig. 2. Surface temperature (°C) curves vs. time (s) during ablation of ZrB2–SiC– ZrO2 ceramics.
found for both ceramics at the beginning of ablation. Hereafter, the temperature reached a steady state with the maximum point as 2070 °C for BS10Z and 2100 °C for BS20Z. Such difference in the surface temperature was probably ascribed to the introduction of ZrO2 with a low thermal conductivity (ZrO2: 2.5 W/(m °C) [25]; ZrB2–SiC: 126.75 W/(m °C) [26]), and hence higher ZrO2 content would bring the lower total thermal conductivity for the ceramic which resulted in the discrepant thermal conductive effect. Meanwhile, the potential measurement error during tests may also lead to the difference in temperature curves of ceramics BS10Z and BS20Z as shown in Fig. 2. Fig. 3 shows the macroscopical photographs of the surface for ZrB2–SiC–ZrO2 ceramics. In the case of ceramic BS10Z, the surface appeared to be covered by a thin white oxidized layer. EDS spectra detected that such oxidized layer was composed of elements of B–
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O–Mg–Si–Zr (Fig. 4). As developed previously [16], ZrB2 got to be oxidized at 800 °C with products of B2O3 and ZrO2, and SiC at 1100 °C with SiO2 and CO. B2O3 and SiO2 tended to form borosilicate at the oxidation environment. Therefore, the oxidized covering the surface of BS10Z after ablation for 30 s was borosilicate. It should be noted that there was some trace of Mg in the oxidized layer, which was attributed to the purity of initial powder of ZrB2 [18]. After ablation for 360 s, the oxidized layer covering the surface of BS10S became apparently denser and thicker than the one of 30 s, which indicated that the ceramic BS10Z underwent serious oxidation. EDS spectra in Fig. 4 demonstrated that peaks of Si and Mg became more intense than the ones of 30 s, which proved evidently that more ZrB2 and SiC in the ceramic BS10Z encountered oxidation under this ablation condition. In addition, some bubbles could be observed from the surface of BS10Z after ablation for 360 s, which was resulted from the high vapor pressure of borosilicate at severe oxidation condition. When the ablation time was extended up to 600 s, distinct oxidized layer was found covering the surface of BS10Z, and the high vapor pressure led to the borosilicate evaporated to form some pores. Concerning the surface of ceramic BS20Z, as shown in Fig. 3, evident discrepancy could be found in comparison with BS10Z. After ablation for 30 s, obvious pores and spallation were observed on the surface of BS20Z. It was more serious that the outer oxidized layer tended to be unripped after ablation with the time exceeded 120 s and then spallated along cooling. That is because of the relaxation of residual stress in the outer oxidized layer during the cooling process after ablation, which meant that the outer oxidized layer covering on the surface of BS20Z did not adhered firmly with the inner part due to the thermal stresses. XRD pattern further revealed that such spallated oxidized layer from ceramic BS20Z was dominantly ZrO2 with monoclinic as the main phase (Fig. 5). 3.3. Microstructure Fig. 6 shows the surface microstructure of ceramic BS10Z after ablation for 30 s. Some bubbles appeared on the surface due to
Fig. 3. Photographs of the surface for BS10Z and BS20Z ceramics after ablation.
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Fig. 4. EDS spectra of ceramic BS10Z after ablation for 30 s and 360 s.
the oxidized reaction beneath the outermost scale during ablation [27], and EDS analysis indicated the bubbles were composed of ZrO2 and borosilicate glass. In the case of BS20Z under the same ablation condition, as shown in Fig. 7, evident difference could be observed from the surface compared with BS10Z. Similar bubbles still appeared but the amounts increased significantly, which meant ceramic BS20Z underwent the severe oxidation during ablation. Moreover, some pores were found on the surface of BS20Z after ablation for 30 s, which was the other difference from BS10Z. Such pores appearing on the surface of BS20Z provided more paths for the inward flow of oxygen [28], and therefore led to the severe oxidation compared with BS10Z. With the extension of ablation time up to 600 s, serious oxidation occurred for ceramic BS10Z as shown in Fig. 8. There were many big conglomerations with some small bubbles dispersed on the surface. High magnified observation revealed that big conglomerations were just the skeleton of ZrO2 produced from the oxidation of ZrB2, and such skeleton was filled with many small pores which allowed for oxygen flow through the oxidized layer during ablation [18,28]. Fig. 9 shows the microstructure of BS20Z after ablation for 600 s with the outermost oxidized layer spallated. It could be found that outer section (layer-1) exhibited severe oxidation than the inner (layer-2), which implied the oxygen flow was effective along both of the longitudinal and transverse directions. Fig. 9b with high
Fig. 5. XRD of the oxidized layer for ceramic BS20Z after ablation for 120 s.
magnification revealed that the oxidized layer in the outer section was dominantly ZrO2, and the inner layer was composed of the mixture of ZrO2 and borosilicate glass. Fig. 9c shows that evident directional growth had been achieved for ZrO2 grains. And obvious pores were further observed on the inner oxidized scale from Fig. 9d, which was favorable for the inward transport of oxygen and led to the serious oxidation for BS20Z during ablation. With regard to the microstructure on the cross-section, Fig. 10 shows the SEM images for the two ceramics after ablation for 30 s. Apparent layered structure was found in the cross-section for both of BS10Z and BS20Z. As denoted by the dotted line, thicker oxidized layer was provided from BS20Z than BS10Z, and BS20Z was more likely to exhibit several layers in the cross-section (BS10Z: two layers; BS20Z: four layers). Total thickness of the oxidized scale was 33 lm for BS10Z and 79 lm for BS20Z. As analyzed above, extensive time for ablation caused the serious oxidized microstructure on the surface layer. Analogously, Fig. 11 shows the microstructure from the cross-section of ceramic BS10Z after ablation for 120 s, which presented evident layered structure for the oxidized scale. The first, i.e., the outermost, was the ZrO2 layer with grains of ZrO2 in columnar shape, which was coincident with the direction of oxygen flow during ablation. The second was the SiC-partly depleted layer with some pores, similarly to the structure of ZrB2–SiC ceramics oxidized at 1900 °C [29]. And the third, i.e., the innermost, was the unreacted material. It could be further indicated that the transition between the first and second layers was composed of small ZrO2 grains and borosilicate glass, besides, the one between the second and third layers was the mixture of borosilicate, ZrO2 grains and unaffected ZrB2 grains. Such layered structure along the cross-section could be testified by the line scanning with EDS in Fig. 11d. Variation in the intensity of different elements represented the detailed constitution in the oxidized layer. Compared with Fig. 10a of BS10Z after ablation 30 s, Fig. 11 illustrates that the oxidized layer with a thickness of 118 lm was obtained for BS10Z after ablation for 120 s. At the same condition, ceramic BS20Z with the outermost oxidized layer spallated (Fig. 3) still remained the layered structure along the cross-section as shown in Fig. 12. The first was the SiC-partly depleted layer mixed with ZrO2 grains and borosilicate glass, and the second was the unaffected layer. Similarly, there were still some pores in the SiC-partly depleted layer. Here two kinds of ZrO2 should be noted as denoted in Figs. 11 and 12. One kind was the ZrO2 in the original material, which distributed uniformly with a mean size of 2 lm. The other is the ZrO2 from the oxidation of ZrB2, which had a smaller size below 1 lm. Such difference in grain size of the two kinds of ZrO2 could be easily found in the cross-sections of ceramics BS10Z and BS20Z.
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Fig. 6. SEM images of surfaces for ceramic BS10Z after ablation for 30 s.
Fig. 7. SEM images of surfaces for ceramic BS20Z after ablation for 30 s: (a) low magnification and (b) high magnification.
Fig. 8. SEM images of surfaces for ceramic BS10Z after ablation for 600 s: (a) low magnification and (b) high magnification.
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Fig. 9. SEM images of surfaces for ceramic BS20Z after ablation for 600 s.
Fig. 10. SEM images of the cross-section for ceramics after ablation for 30 s: (a) BS10Z and (b) BS20Z.
Fig. 13 plots the total thickness of oxidized layer for ZrB2–SiC– ZrO2 ceramics after ablation. Here the thickness was measured along the cross-section direction including the total oxidized scale above the unaffected material. It was found that both of BS10Z and BS20Z provided parabolic-law variation in the thickness of oxidized layer with the extension of ablation time, which was in accordance with the typical oxidation kinetics [16,17]. Compared with BS20Z, ceramic BS10Z possessed thinner oxidized layer at the same condition. And such discrepancy became more obvious with the extended ablation time. In the case of 30 s, Fig. 10 depicts that the thickness of oxidized layer was 33 lm for BS10Z and 79 lm for BS20Z. Regarding the ablation time of 600 s, the total thickness of BS10Z was 164 lm, however, the data for BS20Z was sharply increased up to 749 lm. It was evident that oxidation occurred for BS20Z became more and more severious with the extended ablation time especially compared with BS10Z.
3.4. Weight loss and line gain Fig. 1 has showed the comparison of oxidation behavior of the two ZrB2–SiC–ZrO2 ceramics, and the oxidation during ablation was further discrepant for the two ceramics. Under the ablation condition, products of B2O3 and SiO2 by the oxidation of ZrB2 and SiC would be combined to form borosilicate (Reaction (6)), which was effective to slow down the evaporation of B2O3 at high temperature above 1100 °C [21]. Therefore, resistance to the oxidation of ZrB2–SiC–ZrO2 at that stage came from the restraint of oxygen diffusion by the borosilicate as demonstrated in the part of microstructure. In the case of the ablation for the present ZrB2–SiC–ZrO2 ceramics, the surface temperature was 2000 °C as given in Fig. 1 and Table 1, where the actual oxidation course was quite complicated [19]. Oxidation resistance of ZrB2-based ceramics was achieved by the formation of silica glass layer with
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Fig. 11. SEM images of the cross-section for ceramic BS10Z after ablation for 120 s.
Fig. 12. SEM images of the cross-section for ceramic BS20Z after ablation for 120 s with the surface oxidized layer flaked.
low oxygen permeability produced by the oxidation of SiC [19], which provided an effective protection for the present ZrB2–SiC– ZrO2 ceramics in the same way. Since the surface temperature under the ablation especially for the long exposure time was above 2000 °C and even up to 2100 °C, the effective protection from SiC was obviously weakened due to the active oxidation of SiC according to Reactions (7) and (8), as well as the rapid evaporation of SiO2 via the high vapor pressure (Reaction (9)) [16,17,19].
B2 O3 (l) + SiO2 (l) = Boronsilicate(l)
ð6Þ
SiC(s) + O2 (g) = SiO(g) + CO(g)
ð7Þ
SiC(s) + 2SiO2 (l) = 3SiO(g) + CO(g)
ð8Þ
SiO2 (l) = SiO2 (g)
ð9Þ
During the ablation using oxyacetylene torch, the present ZrB2– SiC–ZrO2 ceramics were also subjected to the large aerodynamic loads in addition to the aerothermal. Therefore, the mechanical loads certainly played an important part during ablation [21]. Under the severe ablation environment, an oxide skeleton structure was formed by the product of ZrO2, which remained configurational stability for the ceramics. Such skeleton combined with SiO2 was responsible for the continuous configurational stability under the conditions coupled with mechanical loads and oxidation, which was explained coincidently from ceramic BS10Z under ablation. In fact, with the ZrO2 content increased from 10 vol% to 20 vol%, the morphology showed evident transformation as analyzed above. Moreover, weight change during ablation also revealed that discrepant ability to resist ablation of the two ceramics. As depicted in Fig. 14, the specific weight loss, calculated according to Eq. (1) presented increment with the extended abla-
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Fig. 13. Thickness (lm) of oxidation layer for BS10Z and BS20Z ceramics after ablation.
Fig. 14. Specific weight loss (mg/cm2) of BS10Z and BS20Z ceramics after ablation.
tion time for both ceramics, and approximate magnitude could be obtained for the short ablation time. However, the difference in the weight loss of the two ceramics became more and more obvious with the increased ablation time. Particularly, the data at 600 s reached 0.01203 mg/cm2 for BS10Z, and the one for BS20Z increased up to about twice of BS10Z, i.e., 0.02065 mg/cm2. In addition to the weight change to describe the oxidation and ablation behavior, reported work has also conducted the height change to further characterize the linear dimension change of the specimens [22]. Because the oxidation during ablation would increase the linear dimension of the specimen, line gain was introduced to compare the ablation behavior of the two ZrB2–SiC– ZrO2 ceramics. Regarding the line gain of the two ceramics calculated Eq. (2), analogous difference was further found in Fig. 15. In the beginning of ablation, both ceramics exhibited a special moment, called ‘‘steady state”, when the line gain almost did not present any variation. During the actual ablation, two conditions were inflicted together on ZrB2–SiC–ZrO2 ceramics, including mechanical scouring and oxidation. Mechanical load caused spallation of the surface layer, and oxidation produced the new oxidized layer at the same time. The two effects contributed in combination for ZrB2–SiC–ZrO2 ceramics and led to the variation in line gain as plotted in Fig. 15. Apparently, the ‘‘steady state” for line gain was
Fig. 15. Specific line gain (mm/s) BS10Z and BS20Z ceramics after ablation.
the approximate equality of mechanical scouring-induced line loss and oxidation-induced line gain. Once the increment in oxidationinduced line gain exceeded the line loss induced by mechanical scouring, the total line gain then presented gradual increment, which could be easily found in Fig. 15. It should be further noted that ceramic BS10Z provided a longer ‘‘steady state” (150 s) during ablation than the one of BS20Z (90 s). Moreover, the line gain of BS20Z became more serious than BS10Z with the extended ablation time. After ablation for 600 s, the specific line gain of BS20Z was up to 4.0 104 mm/s, although the one of BS10Z was only 2.8 104 mm/s. Complementarily, Tang et al. reported that addition of ZrB2–SiC to carbon composites could obviously reduce the thickness erosion rate under the arc-heated wind tunnel ablation [22], and such degradation in line ablation rate could be also found in ZrB2-containing carbon composites under the oxyacetylene torch by Li et al. [30]. Reduction in line ablation rate indicated that sufficient oxidation-induced line gain exceeded the ablation-induced line loss. With respect to the present ZrB2–SiC–ZrO2 ceramics, the line gain generally showed increase along the ablation process especially after the initial ‘‘steady state”, which implied that oxidation was the dominant process during the actual ablation for the present ceramics. Accordingly, oxidation with the dependence on ZrO2 content should be critical to carry out. 3.5. Ablation mechanism Analysis about the microstructure and change in weight and line indicated that BS10Z provided lighter oxidation and ablation than BS20Z, which suggested BS10Z with better oxidation and ablation resistance. According to our previous development [14,15], introduction of ZrO2 was evidently favorable to toughen ZrB2-based ceramics. The present work, however, revealed that excessive ZrO2 was adverse to the ablation and oxidation resistance of ZrB2–SiC–ZrO2 ceramics. Concerning ZrB2–SiC ceramics, reported progress has indicated that different oxidized reaction, with products covering the surface, inhibited the inward transport of oxygen, which was the intrinsic mechanism to achieve oxidation resistance [16]. Although more fraction of ZrO2 was beneficial to toughen the ZrB2-based ceramics [14,15], the essential non-oxidation property of ZrO2 may lead to the probability of lower oxidation resistance for ZrB2–SiC–ZrO2 ceramics. It was well known that oxides produced from the oxidation of ZrB2 and SiC were able to fill and close the microdefects and then resisted the oxygen flow [16]. Nevertheless, microdefects among ZrO2 grains would become the unneglectable path for the inward
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flow of oxygen and then induced the acute oxidation for the matrix material, i.e., ZrB2–SiC. Moreover, despite the benefit induced by addition of ZrO2 to the densification [14,15], increased grain boundaries in ZrB2-based composites, along with the grain size degraded, would be also presented as the potential path for the inward transport of oxygen during ablation. With the increase in ZrO2 content, such assistance to oxygen transport became more evidently due to the serious degraded grain size, i.e., in ceramic BS20Z, which resulted in severe oxidation and ablation for BS20Z. As concluded, addition of ZrO2 was advantageous to the toughening of ZrB2-based ceramics, however, excessive ZrO2 was evidently adverse to the oxidation resistance and further ablation resistance. Consequently in perspectives combined with toughening, oxidation as well as ablation, the addition of ZrO2 to ZrB2-based ceramics should be controlled appropriately. 4. Conclusions By means of the oxyacetylene torch, ablation behavior of ZrB2– SiC–ZrO2 ceramics was investigated in terms of the effect of ZrO2 content. Results of microstructure showed that ceramic BS10Z provided the lighter oxidized structure than BS20Z. Particularly, total thickness of the oxidized layer was 164 lm for BS10Z and 749 lm for BS20Z after ablation for 600 s. TGA further indicated that BS10Z showed initial weight change at higher temperature. Analogously, ceramic BS10Z got a specific weight loss of 0.01203 mg/cm2 after ablation for 600 s, whereas BS20Z of 0.02065 mg/cm2. Regarding the line gain, the two ceramics both exhibited a ‘‘constant line state” at the initial moment of ablation, however, the specific line gain of BS10Z was only half of the one of BS20Z after ablation for 600 s. Ablation mechanism revealed that excessive ZrO2 would supply much path to the inward transport of oxygen flow, which led to the ceramic BS20Z with lower resistance to oxidation and ablation. Acknowledgements This work was supported by the National Natural Science Foundation of China (Nos. 50602010 and 50702016), the National Natural Science Funds for Distinguished Young Scholar (No. 10725207). The authors are also grateful to the anonymous reviewers who provided many valuable comments for improving the presentation and quality of the final paper. References [1] W.G. Fahrenholtz, G.E. Hilmas, I.G. Talmy, J.A. Zaykoski, Refractory diborides of zirconium and hafnium, J. Am. Ceram. Soc. 90 (2007) 1347–1364. [2] J.J. Schuldies, T. Branch, Ultrasonic nde of ceramic components, MCIC Rep. 3 (1978) 429–448. [3] K. Upadhya, J.M. Yang, W.P. Hoffmann, Materials for ultrahigh temperature structural applications, Am. Ceram. Soc. Bull. 76 (1997) 51–56. [4] F. Monteverde, Beneficial effects of an ultra-fine a-SiC incorporation on the sinterability and mechanical properties of ZrB2, Appl. Phys. A 82 (2006) 329– 337.
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[5] A. Rezaie, W.G. Fahrenholtz, G.E. Hilmas, Effect of hot pressing time and temperature on the microstructure and mechanical properties of ZrB2–SiC, J. Mater. Sci. 42 (2007) 2735–2744. [6] A.L. Chamberlain, W.G. Fahrenholtz, G.E. Hilmas, D.T. Ellerby, High-strength zirconium diboride-based ceramics, J. Am. Ceram. Soc. 87 (2004) 1170–1172. [7] F. Monteverde, A. Bellosi, Development and characterization of metal diboridebased composites toughened with ultra-fine SiC particulates, Solid State Sci. 7 (2005) 622–630. [8] F. Monteverde, L. Scatteia, Resistance to thermal shock and to oxidation of metal diborides–SiC ceramics for aerospace application, J. Am. Ceram. Soc. 90 (2007) 1130–1138. [9] M.M. Opeka, I. Talmy, J. Zaykoski, Oxidation-based materials selection for 2000 °C+ hypersonic aerosurfaces: theoretical considerations and historical experience, J. Mater. Soc. 39 (2004) 5887–5904. [10] E.J. Opila, S. Levine, J. Lorincz, Oxidation of ZrB2- and HfB2-based ultra-high temperature ceramics: effects of Ta additions, J. Mater. Sci. 39 (2004) 5969– 5977. [11] I.G. Talmy, J.A. Zaykoski, M.M. Opeka, S. Dallek, Oxidation of ZrB2 ceramics modified with SiC and group IV–VI transition metal diborides, Elec. Chem. Soc. Proc. 12 (2001) 144–158. [12] D. Sciti, M. Brach, A. Bellosi, Oxidation behavior of a pressureless sintered ZrB2–MoSi2 ceramic composites, J. Mater. Res. 20 (2005) 922–930. [13] F. Monteverde, The thermal stability in air of hot-pressed diboride matrix composites for uses at ultra-high temperatures, Corros. Sci. 47 (2005) 2020– 2033. [14] X.H. Zhang, W.J. Li, C.Q. Hong, W.B. Han, Microstructure and mechanical properties of ZrB2-based composites reinforced and toughened by zirconia, Int. J. Appl. Ceram. Technol. 5 (2008) 499–504. [15] W.J. Li, X.H. Zhang, C.Q. Hong, W.B. Han, J.C. Han, Microstructure and mechanical properties of zirconia-toughened ZrB2–MoSi2 composites prepared by hot-pressing, Scripta Mater. 60 (2009) 100–103. [16] F. Monteverde, A. Bellosi, Oxidation of ZrB2-based ceramics in dry air, J. Electrochem. Soc. 150 (2003) B552–B559. [17] A. Rezaie, W.G. Fahrenholtz, G.E. Hilmas, Evolution of structure during the oxidation of zirconium diboride–silicon carbide in air up to 1500 °C, J. Eur. Ceram. Soc. 27 (2007) 2495–2501. [18] S.S. Hwang, A.L. Vasiliev, N.P. Padture, Improved processing and oxidationresistance of ZrB2 ultra-high temperature ceramics containing SiC nanodispersoids, Mater. Sci. Eng. A 464 (2007) 216–224. [19] J.C. Han, P. Hu, X.H. Zhang, S.H. Meng, W.B. Han, Oxidation-resistant ZrB2–SiC composites at 2200 °C, Comp. Sci. Technol. 68 (2008) 799–806. [20] F. Monteverde, R. Savino, Stability of ultra-high-temperature ZrB2–SiC ceramics under simulated atmospheric re-entry conditions, J. Eur. Ceram. Soc. 27 (2007) 4797–4805. [21] X.H. Zhang, P. Hu, J.C. Han, S.H. Meng, Ablation behavior of ZrB2–SiC ultra high temperature ceramics under simulated atmospheric re-entry conditions, Comp. Sci. Technol. 68 (2008) 1718–1726. [22] S.F. Tang, J.Y. Deng, S.J. Wang, W.C. Liu, Comparison of thermal and ablation behaviors of C/SiC composites and C/ZrB2–SiC composites, Corros. Sci. 51 (2009) 54–61. [23] X.H. Zhang, P. Hu, J.C. Han, Structure evolution of ZrB2–SiC during the oxidation in air, J. Mater. Res. 23 (2008) 1961–1972. [24] S. Nakayama, S. Maekawa, T. Sato, Y. Masuda, S. Imai, M. Sakamoto, Mechanical properties of ytterbia stabilized zirconia ceramics (Yb-TZP) fabricated from powders prepared by co-precipitation method, Ceram. Int. 26 (2000) 207–211. [25] R.H.J. Hannink, P.M. Kelly, B.C. Muddle, Transformation toughening in zirconia-containing ceramics, J. Am. Ceram. Soc. 83 (2000) 461–487. [26] W.J. Li, Y. Zhang, X.H. Zhang, C.Q. Hong, W.B. Han, Thermal shock behavior of ZrB2–SiC ultra-high temperature ceramics with addition of zirconia, J. Alloy. Compd. 478 (2009) 386–391. [27] M. Gasch, D. Ellerby, E. Irby, S. Beckman, M. Gusman, S. Johnson, Processing, properties and arc jet oxidation of hafnium/silicon carbide ultra high temperature ceramic, J. Mater. Sci. 39 (2004) 5925–5937. [28] J.C. Han, P. Hu, X.H. Zhang, S.H. Meng, Oxidation behavior of zirconium diboride–silicon carbide at 1800 °C, Scripta Mater. 57 (2007) 825–828. [29] W.B. Han, P. Hu, X.H. Zhang, J.C. Han, S.H. Meng, High-temperature oxidation at 1900 °C of ZrB2–xSiC ultra high temperature ceramic composites, J. Am. Ceram. Soc. 91 (2008) 3328–3334. [30] X.T. Li, J.L. Shi, G.B. Zhang, H. Zhang, Q.G. Guo, L. Liu, Effect of ZrB2 on the ablation properties of carbon composites, Mater. Lett. 60 (2006) 892–896.
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Corrosion Science journal homepage: www.elsevier.com/locate/corsci
Ablation behavior of ZrB2–SiC–ZrO2 ceramic composites by means of the oxyacetylene torch Shanbao Zhou a, Weijie Li a,*, Ping Hu a, Changqing Hong a, Ling Weng b a b
Center for Composite Materials, Harbin Institute of Technology, Harbin 150001, PR China School of Material Science and Engineering, Harbin University of Science and Technology, Harbin 150001, PR China
a r t i c l e
i n f o
Article history: Received 9 March 2009 Accepted 27 May 2009 Available online 3 June 2009 Keywords: A. Ceramic matrix composites B. SEM B. Weight loss C. Oxidation
a b s t r a c t Ablation behavior of ZrB2–SiC–ZrO2 ceramics with two ZrO2 contents was investigated using oxyacetylene torch. Thermogravimetric analysis demonstrated that ceramic with 10 vol% ZrO2 showed initial weight change at higher temperature than the one with 20 vol% ZrO2. After same ablation condition, lighter oxidized microstructure and lower weight loss and line gain were obtained from ceramic with 10 vol% ZrO2. Ablation mechanism revealed that excessive ZrO2 would supply much path to the inward transport of oxygen, which led to the dissatisfactory resistance to oxidation and ablation for the ceramic with 20 vol% ZrO2. Ó 2009 Elsevier Ltd. All rights reserved.
1. Introduction Transition metal borides including ZrB2, HfB2 and TiB2 are of great interest for their unique properties combined with ultra-high melting temperature (>3000 °C), high electrical and thermal conductivities, as well as excellent chemical inertness [1]. Such properties make them the potential candidates for leading edges on aircraft and re-entry vehicles, and for structural parts in high temperature environments [2,3]. Development of these ceramics has focused on densification [4,5], mechanical properties [5–7], thermal shock and oxidation behavior [8–13]. Monolithic ZrB2 was limited in the sinterability due to its covalent bonding, high melting temperature, and low self-diffusion coefficients of Zr and B. In this perspective, SiC was introduced to overcome this obstacle, and near theoretical density was achieved from ZrB2–SiC system. In addition, dissatisfactory toughness of ZrB2 ceramics was still a shortcoming to a wider range of use for most sophisticated applications [1]. Given that defect, phase transformation of ZrO2 was proposed to toughen ZrB2-based ceramics in the previous work [14,15]. As a critical enhancement, oxidation and ablation behavior was investigated widely by many researchers. In terms of the oxidation of ZrB2-based ceramics, ZrB2 is first oxidized to produce a scale composed of ZrO2 and B2O3 at 800 °C under the oxidizing conditions [16]. Due to the lower melting temperature (450 °C), liquid B2O3 can form a continuous layer that seals the surface and acts as an effective barrier to the inward transport of oxygen below * Corresponding author. Tel./fax: +86 451 86402382. E-mail address: [email protected] (W. Li). 0010-938X/$ - see front matter Ó 2009 Elsevier Ltd. All rights reserved. doi:10.1016/j.corsci.2009.05.035
1100 °C. At temperatures above 1100 °C, B2O3 begins to evaporate and such behavior becomes dominant when temperature is higher than 1400 °C due to the high vapor pressure of B2O3 [16–18]. However, the oxidation of ZrB2 can be improved by adding SiC to promote the formation of a silica-rich scale on the exposed surfaces with temperature above 1100 °C. That silica-rich layer is also refractory and then provides resistance to oxygen flow [9,16–18]. Moreover, such layer will resist oxidation effectively up to at least 1500 °C, because SiO2 is evidently less volatile than B2O3 at this temperature [17]. Once the temperature is above 2000 °C and even up to 2100 °C, the effective protection from SiC is obviously weakened due to the active oxidation of SiC to produce SiO, as well as the rapid evaporation of SiO2 via the high vapor pressure [9,19]. ZrO2 skeleton is then dominantly responsible for the oxidation resistance of ZrB2-based ceramics at this stage [9,19]. Under the ablation environment, mechanical loads are inflicted on the ceramics in addition to the oxidation. Much progress has been also carried out in ZrB2-based and ZrB2-conatining ceramics by means of the oxyacetylene torch and arc-heated wind tunnel tests [20–22]. Because of the mechanical loads-induced scouring, weight of the ceramics showed great variation compared with the sole oxidation process [20,22]. Analogously, microstructure evolution was rather complicated for ceramics after ablation, which provided typical layered microstructure along the cross-section [23], as well as the trace of mechanical scouring in ceramics [20–22]. Based on the reported work about ZrO2-toughened ZrB2 ceramics, the present paper was newly proposed to evaluate the ablation behavior of ZrB2–SiC–ZrO2 ceramics using oxyacetylene torch. Effect of ZrO2 content on the ablation behavior was determined in
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2. Experimental
where m1 was the weight (mg) of ceramic specimen before ablation, m2 was the one after ablation, A was the surface area (cm2) of the specimen exposed for ablation. The specific line gain (mm/s) was calculated by measuring the line before and after ablation with the following equation:
2.1. Starting powders and preparation
Dl ¼
Commercially available raw materials were used to prepare the ZrB2–SiC–ZrO2 composite in this work. The ZrB2 powder with a mean size of 2 lm was supplied from the Northwest Institute for Non-Ferrous Metal Research, China. The SiC particles (1 lm, Weifang Kaihua Micro-Powder Co., Ltd., Shandong, China) were predominantly a-SiC. The ZrO2 (1 lm, Fanmeiya Powders Co., Ltd., Jiangxi, China) used here was 3 mol% Y2O3 partially stabilized zirconia prepared by the co-precipitation method [24]. Two fractions of ZrO2, i.e., 10 vol% and 20 vol%, were considered to add to ZrB2–10 vol%SiC ceramics, and the two ceramics were marked as BS10Z and BS20Z, respectively. According to the detailed composition, powder mixtures were milled for 8 h in a polyethylene bottle with ZrO2 balls and ethanol as the grinding media. After mixing, the slurry was dried in a rotary evaporator and screened. Milled powders were hot-pressed at 1850 °C for 60 min under a uniaxial load of 30 MPa in Ar atmosphere.
where l1 and l2 were the height (mm) of the specimen before and after ablation, t was the ablation time (s).
terms of the ablation and oxidation mechanism with investigating the microstructure as well as the variation of weight and line after ablation.
2.2. Ablation tests Ablation behavior of the present ZrB2–SiC–ZrO2 ceramics was evaluated by means of the oxyacetylene torch. Table 1 lists the specific experimental condition and the measured surface temperature. During the ablation tests, a specimen with 15 mm in diameter and 10 mm in height was exposed to the flame. Ablation time was chosen as 30 s, 60 s, 120 s, 180 s, 240 s, 360 s and 600 s. The distance between the nozzle tip of the oxyacetylene gun and the surface of ceramic specimen was 10 mm. In order to ensure the ablation in just one direction, the specimen was held in a concave graphite anvil and only one face was exposed to the ablation environment. Flow rates of the oxygen and acetylene were 1.02 L/s and 0.36 L/s, respectively. The inner diameter of the nozzle tip of the ablation gun was 2 mm. Surface temperature of the ceramic samples was determined by means of an optical pyrometer and recorded every 10 s. 2.3. Characterization Crystalline phases of the specimen after ablation were identified by X-ray diffraction (XRD, Rigaku, Dmax-rb). The microstructure after ablation was analyzed by scanning electron microscopy/energy dispersive spectroscopy (SEM/EDS, Quanta 200). Thermogravimetric analysis (TGA) was conducted on a Netzsch STA 449C using a heating rate of 10 °C/min from 20 °C to 1400 °C. The specific weight loss (mg/cm2) was calculated by measuring the weight before and after ablation with the following equation:
Dm ¼
m1 m2 A
l2 l1 t
ð2Þ
3. Results and discussion 3.1. TGA TGA experiments were conducted firstly for the two ZrB2–SiC– ZrO2 ceramics from 20 °C (room temperature) to 1400 °C to assist exploring the ablation mechanism. As plotted in Fig. 1, the weight did not change with the increased testing temperature for both ceramics below 700 °C. The first time for weight change of ceramic BS10Z occurred at 795 °C, and BS20Z at 735 °C. Such weight change indicated the presence of the initial oxidation of ZrB2 in ZrB2–SiC–ZrO2 ceramics according to Reaction (3). As the temperature elevated, weight change increased gradually for both ceramics. Previous studies reported that B2O3 evaporated rapidly (Reaction (4)) at temperature above 1100 °C [16–19]. At 1150 °C in Fig. 1, an evident inflection point appeared for the weight change curve of ceramic BS20Z, and similar inflection point for BS10Z was at 1220 °C. Such point represented that SiC in ZrB2–SiC–ZrO2 ceramics began to be oxidized rapidly according to Reaction (5), which was basically consistent with the reported progress [16]. After that, the weight for both ceramics both provided monotonous increment with extended temperature.
ZrB2 (s) + 5/2O2 (g) = ZrO2 (s) + B2 O3 (l)
ð3Þ
B2 O3 (l) = B2 O3 (g)
ð4Þ
SiC(s) + 3/2O2 (g) = SiO2 (l) + CO(g)
ð5Þ
3.2. Macroscopical morphology Fig. 2 plots the temperature curves on the surfaces of the two ZrB2–SiC–ZrO2 ceramics varied with the ablation time increased up to 600 s. Sharp increment in the surface temperature could be
ð1Þ
Table 1 Experimental condition for oxyacetylene ablation and the measured surface temperature. O2 gas pressure (kPa) O2 gas flux (L/s) C2H2 gas pressure (kPa) C2H2 gas flux (L/s) Diameter of nozzle (mm) Distance from sample surface to nozzle (mm) Measured surface temperature (°c)
0.45 1.02 0.1 0.36 2 10 2000 ± 1
Fig. 1. TGA of BS10Z and BS20Z ceramics in air.
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Fig. 2. Surface temperature (°C) curves vs. time (s) during ablation of ZrB2–SiC– ZrO2 ceramics.
found for both ceramics at the beginning of ablation. Hereafter, the temperature reached a steady state with the maximum point as 2070 °C for BS10Z and 2100 °C for BS20Z. Such difference in the surface temperature was probably ascribed to the introduction of ZrO2 with a low thermal conductivity (ZrO2: 2.5 W/(m °C) [25]; ZrB2–SiC: 126.75 W/(m °C) [26]), and hence higher ZrO2 content would bring the lower total thermal conductivity for the ceramic which resulted in the discrepant thermal conductive effect. Meanwhile, the potential measurement error during tests may also lead to the difference in temperature curves of ceramics BS10Z and BS20Z as shown in Fig. 2. Fig. 3 shows the macroscopical photographs of the surface for ZrB2–SiC–ZrO2 ceramics. In the case of ceramic BS10Z, the surface appeared to be covered by a thin white oxidized layer. EDS spectra detected that such oxidized layer was composed of elements of B–
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O–Mg–Si–Zr (Fig. 4). As developed previously [16], ZrB2 got to be oxidized at 800 °C with products of B2O3 and ZrO2, and SiC at 1100 °C with SiO2 and CO. B2O3 and SiO2 tended to form borosilicate at the oxidation environment. Therefore, the oxidized covering the surface of BS10Z after ablation for 30 s was borosilicate. It should be noted that there was some trace of Mg in the oxidized layer, which was attributed to the purity of initial powder of ZrB2 [18]. After ablation for 360 s, the oxidized layer covering the surface of BS10S became apparently denser and thicker than the one of 30 s, which indicated that the ceramic BS10Z underwent serious oxidation. EDS spectra in Fig. 4 demonstrated that peaks of Si and Mg became more intense than the ones of 30 s, which proved evidently that more ZrB2 and SiC in the ceramic BS10Z encountered oxidation under this ablation condition. In addition, some bubbles could be observed from the surface of BS10Z after ablation for 360 s, which was resulted from the high vapor pressure of borosilicate at severe oxidation condition. When the ablation time was extended up to 600 s, distinct oxidized layer was found covering the surface of BS10Z, and the high vapor pressure led to the borosilicate evaporated to form some pores. Concerning the surface of ceramic BS20Z, as shown in Fig. 3, evident discrepancy could be found in comparison with BS10Z. After ablation for 30 s, obvious pores and spallation were observed on the surface of BS20Z. It was more serious that the outer oxidized layer tended to be unripped after ablation with the time exceeded 120 s and then spallated along cooling. That is because of the relaxation of residual stress in the outer oxidized layer during the cooling process after ablation, which meant that the outer oxidized layer covering on the surface of BS20Z did not adhered firmly with the inner part due to the thermal stresses. XRD pattern further revealed that such spallated oxidized layer from ceramic BS20Z was dominantly ZrO2 with monoclinic as the main phase (Fig. 5). 3.3. Microstructure Fig. 6 shows the surface microstructure of ceramic BS10Z after ablation for 30 s. Some bubbles appeared on the surface due to
Fig. 3. Photographs of the surface for BS10Z and BS20Z ceramics after ablation.
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Fig. 4. EDS spectra of ceramic BS10Z after ablation for 30 s and 360 s.
the oxidized reaction beneath the outermost scale during ablation [27], and EDS analysis indicated the bubbles were composed of ZrO2 and borosilicate glass. In the case of BS20Z under the same ablation condition, as shown in Fig. 7, evident difference could be observed from the surface compared with BS10Z. Similar bubbles still appeared but the amounts increased significantly, which meant ceramic BS20Z underwent the severe oxidation during ablation. Moreover, some pores were found on the surface of BS20Z after ablation for 30 s, which was the other difference from BS10Z. Such pores appearing on the surface of BS20Z provided more paths for the inward flow of oxygen [28], and therefore led to the severe oxidation compared with BS10Z. With the extension of ablation time up to 600 s, serious oxidation occurred for ceramic BS10Z as shown in Fig. 8. There were many big conglomerations with some small bubbles dispersed on the surface. High magnified observation revealed that big conglomerations were just the skeleton of ZrO2 produced from the oxidation of ZrB2, and such skeleton was filled with many small pores which allowed for oxygen flow through the oxidized layer during ablation [18,28]. Fig. 9 shows the microstructure of BS20Z after ablation for 600 s with the outermost oxidized layer spallated. It could be found that outer section (layer-1) exhibited severe oxidation than the inner (layer-2), which implied the oxygen flow was effective along both of the longitudinal and transverse directions. Fig. 9b with high
Fig. 5. XRD of the oxidized layer for ceramic BS20Z after ablation for 120 s.
magnification revealed that the oxidized layer in the outer section was dominantly ZrO2, and the inner layer was composed of the mixture of ZrO2 and borosilicate glass. Fig. 9c shows that evident directional growth had been achieved for ZrO2 grains. And obvious pores were further observed on the inner oxidized scale from Fig. 9d, which was favorable for the inward transport of oxygen and led to the serious oxidation for BS20Z during ablation. With regard to the microstructure on the cross-section, Fig. 10 shows the SEM images for the two ceramics after ablation for 30 s. Apparent layered structure was found in the cross-section for both of BS10Z and BS20Z. As denoted by the dotted line, thicker oxidized layer was provided from BS20Z than BS10Z, and BS20Z was more likely to exhibit several layers in the cross-section (BS10Z: two layers; BS20Z: four layers). Total thickness of the oxidized scale was 33 lm for BS10Z and 79 lm for BS20Z. As analyzed above, extensive time for ablation caused the serious oxidized microstructure on the surface layer. Analogously, Fig. 11 shows the microstructure from the cross-section of ceramic BS10Z after ablation for 120 s, which presented evident layered structure for the oxidized scale. The first, i.e., the outermost, was the ZrO2 layer with grains of ZrO2 in columnar shape, which was coincident with the direction of oxygen flow during ablation. The second was the SiC-partly depleted layer with some pores, similarly to the structure of ZrB2–SiC ceramics oxidized at 1900 °C [29]. And the third, i.e., the innermost, was the unreacted material. It could be further indicated that the transition between the first and second layers was composed of small ZrO2 grains and borosilicate glass, besides, the one between the second and third layers was the mixture of borosilicate, ZrO2 grains and unaffected ZrB2 grains. Such layered structure along the cross-section could be testified by the line scanning with EDS in Fig. 11d. Variation in the intensity of different elements represented the detailed constitution in the oxidized layer. Compared with Fig. 10a of BS10Z after ablation 30 s, Fig. 11 illustrates that the oxidized layer with a thickness of 118 lm was obtained for BS10Z after ablation for 120 s. At the same condition, ceramic BS20Z with the outermost oxidized layer spallated (Fig. 3) still remained the layered structure along the cross-section as shown in Fig. 12. The first was the SiC-partly depleted layer mixed with ZrO2 grains and borosilicate glass, and the second was the unaffected layer. Similarly, there were still some pores in the SiC-partly depleted layer. Here two kinds of ZrO2 should be noted as denoted in Figs. 11 and 12. One kind was the ZrO2 in the original material, which distributed uniformly with a mean size of 2 lm. The other is the ZrO2 from the oxidation of ZrB2, which had a smaller size below 1 lm. Such difference in grain size of the two kinds of ZrO2 could be easily found in the cross-sections of ceramics BS10Z and BS20Z.
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Fig. 6. SEM images of surfaces for ceramic BS10Z after ablation for 30 s.
Fig. 7. SEM images of surfaces for ceramic BS20Z after ablation for 30 s: (a) low magnification and (b) high magnification.
Fig. 8. SEM images of surfaces for ceramic BS10Z after ablation for 600 s: (a) low magnification and (b) high magnification.
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Fig. 9. SEM images of surfaces for ceramic BS20Z after ablation for 600 s.
Fig. 10. SEM images of the cross-section for ceramics after ablation for 30 s: (a) BS10Z and (b) BS20Z.
Fig. 13 plots the total thickness of oxidized layer for ZrB2–SiC– ZrO2 ceramics after ablation. Here the thickness was measured along the cross-section direction including the total oxidized scale above the unaffected material. It was found that both of BS10Z and BS20Z provided parabolic-law variation in the thickness of oxidized layer with the extension of ablation time, which was in accordance with the typical oxidation kinetics [16,17]. Compared with BS20Z, ceramic BS10Z possessed thinner oxidized layer at the same condition. And such discrepancy became more obvious with the extended ablation time. In the case of 30 s, Fig. 10 depicts that the thickness of oxidized layer was 33 lm for BS10Z and 79 lm for BS20Z. Regarding the ablation time of 600 s, the total thickness of BS10Z was 164 lm, however, the data for BS20Z was sharply increased up to 749 lm. It was evident that oxidation occurred for BS20Z became more and more severious with the extended ablation time especially compared with BS10Z.
3.4. Weight loss and line gain Fig. 1 has showed the comparison of oxidation behavior of the two ZrB2–SiC–ZrO2 ceramics, and the oxidation during ablation was further discrepant for the two ceramics. Under the ablation condition, products of B2O3 and SiO2 by the oxidation of ZrB2 and SiC would be combined to form borosilicate (Reaction (6)), which was effective to slow down the evaporation of B2O3 at high temperature above 1100 °C [21]. Therefore, resistance to the oxidation of ZrB2–SiC–ZrO2 at that stage came from the restraint of oxygen diffusion by the borosilicate as demonstrated in the part of microstructure. In the case of the ablation for the present ZrB2–SiC–ZrO2 ceramics, the surface temperature was 2000 °C as given in Fig. 1 and Table 1, where the actual oxidation course was quite complicated [19]. Oxidation resistance of ZrB2-based ceramics was achieved by the formation of silica glass layer with
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Fig. 11. SEM images of the cross-section for ceramic BS10Z after ablation for 120 s.
Fig. 12. SEM images of the cross-section for ceramic BS20Z after ablation for 120 s with the surface oxidized layer flaked.
low oxygen permeability produced by the oxidation of SiC [19], which provided an effective protection for the present ZrB2–SiC– ZrO2 ceramics in the same way. Since the surface temperature under the ablation especially for the long exposure time was above 2000 °C and even up to 2100 °C, the effective protection from SiC was obviously weakened due to the active oxidation of SiC according to Reactions (7) and (8), as well as the rapid evaporation of SiO2 via the high vapor pressure (Reaction (9)) [16,17,19].
B2 O3 (l) + SiO2 (l) = Boronsilicate(l)
ð6Þ
SiC(s) + O2 (g) = SiO(g) + CO(g)
ð7Þ
SiC(s) + 2SiO2 (l) = 3SiO(g) + CO(g)
ð8Þ
SiO2 (l) = SiO2 (g)
ð9Þ
During the ablation using oxyacetylene torch, the present ZrB2– SiC–ZrO2 ceramics were also subjected to the large aerodynamic loads in addition to the aerothermal. Therefore, the mechanical loads certainly played an important part during ablation [21]. Under the severe ablation environment, an oxide skeleton structure was formed by the product of ZrO2, which remained configurational stability for the ceramics. Such skeleton combined with SiO2 was responsible for the continuous configurational stability under the conditions coupled with mechanical loads and oxidation, which was explained coincidently from ceramic BS10Z under ablation. In fact, with the ZrO2 content increased from 10 vol% to 20 vol%, the morphology showed evident transformation as analyzed above. Moreover, weight change during ablation also revealed that discrepant ability to resist ablation of the two ceramics. As depicted in Fig. 14, the specific weight loss, calculated according to Eq. (1) presented increment with the extended abla-
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Fig. 13. Thickness (lm) of oxidation layer for BS10Z and BS20Z ceramics after ablation.
Fig. 14. Specific weight loss (mg/cm2) of BS10Z and BS20Z ceramics after ablation.
tion time for both ceramics, and approximate magnitude could be obtained for the short ablation time. However, the difference in the weight loss of the two ceramics became more and more obvious with the increased ablation time. Particularly, the data at 600 s reached 0.01203 mg/cm2 for BS10Z, and the one for BS20Z increased up to about twice of BS10Z, i.e., 0.02065 mg/cm2. In addition to the weight change to describe the oxidation and ablation behavior, reported work has also conducted the height change to further characterize the linear dimension change of the specimens [22]. Because the oxidation during ablation would increase the linear dimension of the specimen, line gain was introduced to compare the ablation behavior of the two ZrB2–SiC– ZrO2 ceramics. Regarding the line gain of the two ceramics calculated Eq. (2), analogous difference was further found in Fig. 15. In the beginning of ablation, both ceramics exhibited a special moment, called ‘‘steady state”, when the line gain almost did not present any variation. During the actual ablation, two conditions were inflicted together on ZrB2–SiC–ZrO2 ceramics, including mechanical scouring and oxidation. Mechanical load caused spallation of the surface layer, and oxidation produced the new oxidized layer at the same time. The two effects contributed in combination for ZrB2–SiC–ZrO2 ceramics and led to the variation in line gain as plotted in Fig. 15. Apparently, the ‘‘steady state” for line gain was
Fig. 15. Specific line gain (mm/s) BS10Z and BS20Z ceramics after ablation.
the approximate equality of mechanical scouring-induced line loss and oxidation-induced line gain. Once the increment in oxidationinduced line gain exceeded the line loss induced by mechanical scouring, the total line gain then presented gradual increment, which could be easily found in Fig. 15. It should be further noted that ceramic BS10Z provided a longer ‘‘steady state” (150 s) during ablation than the one of BS20Z (90 s). Moreover, the line gain of BS20Z became more serious than BS10Z with the extended ablation time. After ablation for 600 s, the specific line gain of BS20Z was up to 4.0 104 mm/s, although the one of BS10Z was only 2.8 104 mm/s. Complementarily, Tang et al. reported that addition of ZrB2–SiC to carbon composites could obviously reduce the thickness erosion rate under the arc-heated wind tunnel ablation [22], and such degradation in line ablation rate could be also found in ZrB2-containing carbon composites under the oxyacetylene torch by Li et al. [30]. Reduction in line ablation rate indicated that sufficient oxidation-induced line gain exceeded the ablation-induced line loss. With respect to the present ZrB2–SiC–ZrO2 ceramics, the line gain generally showed increase along the ablation process especially after the initial ‘‘steady state”, which implied that oxidation was the dominant process during the actual ablation for the present ceramics. Accordingly, oxidation with the dependence on ZrO2 content should be critical to carry out. 3.5. Ablation mechanism Analysis about the microstructure and change in weight and line indicated that BS10Z provided lighter oxidation and ablation than BS20Z, which suggested BS10Z with better oxidation and ablation resistance. According to our previous development [14,15], introduction of ZrO2 was evidently favorable to toughen ZrB2-based ceramics. The present work, however, revealed that excessive ZrO2 was adverse to the ablation and oxidation resistance of ZrB2–SiC–ZrO2 ceramics. Concerning ZrB2–SiC ceramics, reported progress has indicated that different oxidized reaction, with products covering the surface, inhibited the inward transport of oxygen, which was the intrinsic mechanism to achieve oxidation resistance [16]. Although more fraction of ZrO2 was beneficial to toughen the ZrB2-based ceramics [14,15], the essential non-oxidation property of ZrO2 may lead to the probability of lower oxidation resistance for ZrB2–SiC–ZrO2 ceramics. It was well known that oxides produced from the oxidation of ZrB2 and SiC were able to fill and close the microdefects and then resisted the oxygen flow [16]. Nevertheless, microdefects among ZrO2 grains would become the unneglectable path for the inward
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flow of oxygen and then induced the acute oxidation for the matrix material, i.e., ZrB2–SiC. Moreover, despite the benefit induced by addition of ZrO2 to the densification [14,15], increased grain boundaries in ZrB2-based composites, along with the grain size degraded, would be also presented as the potential path for the inward transport of oxygen during ablation. With the increase in ZrO2 content, such assistance to oxygen transport became more evidently due to the serious degraded grain size, i.e., in ceramic BS20Z, which resulted in severe oxidation and ablation for BS20Z. As concluded, addition of ZrO2 was advantageous to the toughening of ZrB2-based ceramics, however, excessive ZrO2 was evidently adverse to the oxidation resistance and further ablation resistance. Consequently in perspectives combined with toughening, oxidation as well as ablation, the addition of ZrO2 to ZrB2-based ceramics should be controlled appropriately. 4. Conclusions By means of the oxyacetylene torch, ablation behavior of ZrB2– SiC–ZrO2 ceramics was investigated in terms of the effect of ZrO2 content. Results of microstructure showed that ceramic BS10Z provided the lighter oxidized structure than BS20Z. Particularly, total thickness of the oxidized layer was 164 lm for BS10Z and 749 lm for BS20Z after ablation for 600 s. TGA further indicated that BS10Z showed initial weight change at higher temperature. Analogously, ceramic BS10Z got a specific weight loss of 0.01203 mg/cm2 after ablation for 600 s, whereas BS20Z of 0.02065 mg/cm2. Regarding the line gain, the two ceramics both exhibited a ‘‘constant line state” at the initial moment of ablation, however, the specific line gain of BS10Z was only half of the one of BS20Z after ablation for 600 s. Ablation mechanism revealed that excessive ZrO2 would supply much path to the inward transport of oxygen flow, which led to the ceramic BS20Z with lower resistance to oxidation and ablation. Acknowledgements This work was supported by the National Natural Science Foundation of China (Nos. 50602010 and 50702016), the National Natural Science Funds for Distinguished Young Scholar (No. 10725207). The authors are also grateful to the anonymous reviewers who provided many valuable comments for improving the presentation and quality of the final paper. References [1] W.G. Fahrenholtz, G.E. Hilmas, I.G. Talmy, J.A. Zaykoski, Refractory diborides of zirconium and hafnium, J. Am. Ceram. Soc. 90 (2007) 1347–1364. [2] J.J. Schuldies, T. Branch, Ultrasonic nde of ceramic components, MCIC Rep. 3 (1978) 429–448. [3] K. Upadhya, J.M. Yang, W.P. Hoffmann, Materials for ultrahigh temperature structural applications, Am. Ceram. Soc. Bull. 76 (1997) 51–56. [4] F. Monteverde, Beneficial effects of an ultra-fine a-SiC incorporation on the sinterability and mechanical properties of ZrB2, Appl. Phys. A 82 (2006) 329– 337.
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