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SiC composite ceramics via carbothermal reduction of aluminosilicate precursor for solar sensible thermal storage
Journal of Alloys and Compounds 662 (2016) 126e137
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Journal of Alloys and Compounds journal homepage: http://www.elsevier.com/locate/jalcom
Synthesis and characterization of Al2O3/SiC composite ceramics via carbothermal reduction of aluminosilicate precursor for solar sensible thermal storage Xiaohong Xu, Xinbin Lao*, Jianfeng Wu, Yaxiang Zhang, Xiaoyang Xu, Kun Li State Key Laboratory of Silicate Materials for Architectures, Wuhan University of Technology, Wuhan 430070, PR China
a r t i c l e i n f o
a b s t r a c t
Article history: Received 25 August 2015 Received in revised form 2 December 2015 Accepted 7 December 2015 Available online 12 December 2015
Al2O3 material with high heat capacity and high thermal conductivity was in situ formed to bond with SiC for synthesizing Al2O3/SiC composite ceramics by the removal of silica in aluminosilicate precursor using carbothermal reduction method. Phase transformation, microstructural evolution, and the relevant variations in properties were studied in the temperature range 1460e1580 C. Results indicated that the addition of Y2O3 improved the carbothermal reaction rate significantly and decreased the finish temperature of carbothermal reaction to 1500 C. Incorporation of the in situ formed Al2O3 could lower the sintering temperature of SiC ceramic without decreasing the heat capacity. By sintering at the optimal temperature as 1540 C, Al2O3/SiC composite ceramics with the heat capacity of 1.16 J/(g$K) and the high thermal conductivity of 13.73 W/(m$K) were obtained. The relatively high thermal conductivity endowed the composites with a good thermal shock resistance. This study was intended to identify the reaction conditions for obtaining sensible thermal storage materials with favorable morphology and properties by using a simple in situ synthesis method. © 2015 Elsevier B.V. All rights reserved.
Keywords: Al2O3/SiC composite ceramics Solar sensible thermal storage Carbothermal reduction Heat capacity Thermal conductivity
1. Introduction CO2-induced global warming has become a pressing issue, and needs to be tackled urgently. Efficient utilization of solar energy is thereby proposed and being considered increasingly as a way to achieve a sustainable and clean energy supply for human beings [1,2]. Solar thermal power generation is one of the prevalent power generation technologies using solar energy, and the conversion of solar energy into thermal energy is the easiest and the most extensively accepted method [3]. The main problem of solar thermal power generation is the mismatch between the solar energy and the energy demand, since solar energy is an intermittent energy resource. To remove the fluctuations caused by the intermittent nature of solar energy, a thermal storage system composed of thermal storage materials is required to be attached with solar collectors to store energy and ensure the continuous power supply [4,5]. The thermal storage materials are classified into three main categories according to different storage technologies: sensible heat storage (SHS), latent heat storage, thermo-chemical storage.
* Corresponding author. E-mail address: [email protected] (X. Lao). http://dx.doi.org/10.1016/j.jallcom.2015.12.038 0925-8388/© 2015 Elsevier B.V. All rights reserved.
Among these thermal storage technologies, sensible heat storage is the most simple and inexpensive way for thermal storage system with technological and economical superiority, although there are few advantages of latent storage and thermo-chemical storage over sensible heat storage [6,7]. Excellent properties of the thermal storage materials are the key factors to guarantee the operation of solar sensible thermal storage system. A high thermal storage capacity (i.e. sensible heat capacity) is essential to reduce the volume of sensible thermal storage system and increase the thermal storage efficiency, whilst a good thermal shock resistance allows the thermal storage materials to avoid mechanical degradation after thermal cycles [2]. Numerous oxide ceramics, such as alumina, zirconia, mullite, cordierite, were discovered to be suitable for SHS due to the high thermal storage capacity, the good thermo-chemical stability and the good corrosion resistance against hot heat transfer fluid (HTF, usually air) [8]. Owing to the poor thermal conductivity, however, oxide ceramics were susceptible to thermal shock, and the drastic degradation in mechanical properties after the thermal shock had limited their wide applications at high temperature (850 C), especially the thermal storage application [9,10]. Also the poor thermal conductivity of the oxide ceramics led to the poor heat transfer capability,
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which might slow down thermal storage-retrieval process of the thermal storage system [11,12]. SiC ceramics consequently attracted much attention as candidate materials for solar sensible thermal storage due to the high thermal conductivity as well as the high thermal storage capacity [13,14]. The high sintering temperature (2100 C) of covalent compound SiC [15,16] and the high cost of SiC starting materials [17], however, rendered the difficulties of producing SiC ceramics with low manufacturing cost. Due to the high sintering temperature, the complex sintering processes of SiC ceramics, such as the hot-pressing sintering [18], the spark plasma sintering [19,20], and the vacuum sintering [21], etc, were generally employed for the production of SiC ceramics, which imposed strict requirement on sintering equipment and made the fabrication of products with complex shapes extremely difficult [22], thus hampering the comprehensive utilization of SiC ceramics in solar sensible thermal storage. As known, pressureless liquid phase sintering of SiC by introducing oxide materials as sintering additives did not suffer from these shortcomings, and the introduced oxide materials (e.g. Al2O3, silica, and mullite, etc.) not only promoted the sintering of SiC, but also substituted partial SiC to synthesize SiC composite ceramics with lower cost [23e26]. Among the commonly-used oxide materials, Al2O3 possesses optimal thermophysical properties. For example, in the temperature range 25e600 C, the heat capacity of Al2O3 is 0.77e1.19 J/(g$K), while those of silica and mullite are 0.44e1.13 J/(g$K), 0.76e1.18 J/(g$K), respectively [27,28]; the thermal conductivity of Al2O3 is 17.5 W/ (m$K) at room temperature, which exceeds the thermal conductivity levels of silica and mullite (i.e. 1.4 W/(m$K) for silica, 6.1 W/ (m$K) for mullite) [29e31]. In our previous works [13,31,32], Al2O3eSiC composite ceramics with the high thermal storage capacity, such as Al2O3eSiC and Al2O3eSiCeZrO2 composites, were obtained by using aluminosilicate minerals (e.g. kaolin, andalusite, talc, etc.) as additives to promote the pressureless liquid phase sintering process, whilst the properties were characterized for elevating the suitability for solar sensible thermal storage. The composite ceramics suffered from the poor thermal shock resistance as the bending strength decreased significantly after 30 cycles of thermal shock (instantaneous temperature up to 1000 C), although the high thermal storage capacity and the low sintering temperature were accomplished. The poor thermal shock resistance of the composites was attributed to the poor thermal conductivity of 1.5e2.3 W/(m$K), which mainly resulted from the introduction of silica from the aluminosilicate additives. Hence there was an urgent need to remove the negative effect of silica on the thermal conductivity when using aluminosilicate minerals of low cost as raw materials for fabricating thermal storage materials. It is noteworthy that, as reported by R. Naghizadeh and F. Golestani-fard [33], silica in the aluminosilicates could be reduced in the form of gaseous silicon monoxide (SiO) by carbothermal reduction, whose principle was described in Eq. (1). This method is thereby considered to be an effective way to cope with the thermal conductivity degradation caused by the introduced silica.
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the composite ceramics were investigated, whilst the relevant variations in properties were studied in detail. The Al2O3/SiC composite ceramics were characterized by the high thermal storage capacity, the high thermal conductivity as well as the good thermal shock resistance. The present work was aimed at studying the reaction conditions for obtaining Al2O3/SiC composite thermal storage ceramics with favorable properties by a less complex and lowcost in situ synthesis method. 2. Material and methods 2.1. Starting materials Starting materials for experiments were SiC powder with two particle sizes of 106 mm and 58 mm (Dan Jiangkou Hongyuan Silicon Carbide Co., Ltd, Hubei, China), andalusite (under 250-mesh, Xinjiang Baoan New Energy-Mineral Co.,Ltd, Xinjiang, China), Al powder (under 250-mesh, Sinopharm Chemical Reagent Co.,Ltd, Shanghai, China) and Y2O3 (chemical pure, Sinopharm Chemical Reagent Co.,Ltd, Shanghai, China). Based on our previous study [31], the weight ratio of SiC to Al2O3 in the final products were taken as 65:35, respectively (i.e. 35 wt.% SiC was substituted by Al2O3). According to Eq. (1), the proportions of the starting materials for formula AS were calculated to be: 63.5 wt.% SiC, 22.7 wt.% andalusite, 13.8 wt.% Al. 3 wt.% Y2O3 was added as the catalyst for the decomposition of mullite as well as the sintering additive for densification, since the finish temperature for carbothermal reaction reported in Ref. [33] was as high as 1650 C. Al powder was added to react with the excess silica of andalusite to produce the intermediate mullite. The chemical compositions of starting materials and formula AS were listed in Table 1, in which the compositions of formula AS were calculated including Al2O3 derived from the transformation of Al. The other two formulae, i.e. AS1 and AS2, were designed to study the effects of Y2O3 and the added SiC on the carbothermal reaction rate and the microstructural evolution by performing XRD and SEM analysis of as-sintered samples. Both formulae excluded the addition of SiC, and the ratio of andalusite:Al was consistent with that of formula AS. Formulae AS1 and AS2 were differentiated by 8 wt.% Y2O3 addition and no Y2O3 addition, respectively. 2.2. Processing The starting materials were mixed thoroughly with 5 wt.% PVA binder addition by ball milling, and then pressed by a uniaxial pressure of 20 MPa into rectangular samples of 6.8 mm 6.8 mm 38 mm and cylindrical samples of F10 2 mm. The formed powder compacts were dried at 100 C for 24 h in an electric oven and pressureless sintered by the carbon-buried sintering method (Fig. 1) from 1440 C to 1580 C in a seggar full of carbon at molybdenum disiliciade furnace with a heating rate of 5 C/min and a holding time of 3 h at the maximum temperatures. 2.3. Characterization
3Al2 O3 $2SiO2 ðsÞ þ 2COðgÞ/3Al2 O3 ðsÞ þ 2SiOðgÞþ2CO2 ðgÞ (1) In this study, the carbothermal reduction method was employed to remove the negative effects of silica on the thermal conductivity as well as the heat capacity of Al2O3/SiC composite ceramics, which were fabricated using andalusite, an abundant aluminosilicate mineral with high alumina content, thus avoiding the conventional costly, energy-consuming and complicated sintering processes of SiC ceramics. The main factors affecting the carbothermal reaction rate, the phase transformation and the microstructural evolution of
The water absorption (Wa), the open porosity (Pa) and the bulk density (D) were measured by with the rectangular samples by AUY120 electronic analytical balance (Japanese Shimazu) through static weighing method. The bending strength was measured with the rectangular samples by Computer Control Electronic Universal Test Machine (Shenzhen Reger Instrument Co., Ltd., Guangdong, China). The heat capacity, the thermal diffusivity as well as the thermal conductivity were measured using the cylindrical samples by TC-7000H laser flash thermal constant analyzer(ULVAC SINKURIKO. Inc., Yokohama, Japan). The thermal shock resistance of the
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Table 1 Chemical compositions of starting materials and formula AS (wt%). Starting material
SiC
SiO2
Al2O3
Fe2O3
TiO2
CaO
MgO
K2O
Na2O
I.L
Total
Silicon carbide Andalusite Formula AS
99.19 0 56.27
0 43.25 8.77
0.17 50.87 33.7
0.19 1.11 0.33
0 0.16 0.03
0.09 1.28 0.31
0 0.14 0.03
0 0.74 0.15
0 0.53 0.11
0 1.49 0.3
99.63 99.57 100.0
Fig. 1. Systematic representation of pressureless carbon-buried sintering.
samples was tested under a thermal cycle described as follows: the samples were placed in a furnace and heated to the preset of temperature (1100 C), soaked for 20 min and then were taken out with cooling by wind (at a speed of 4 m/s) to room temperature (25 C). The bending strength loss rate of the samples after thermal shock tests represented the thermal shock resistance. Reported values were available from the average of at least 5 samples. X-ray diffraction (XRD) patterns were collected using a D/MAXIIIAdiffractometer (Rigaku Corporation, Japan) equipped with Cu Ka (l ¼ 1.54 Å) radiation. The relative content of different phases was calculated by RIR (Reference Intensity Ratio) method semiquantitatively [34]. The morphology was observed by SEM (JEOL Ltd., Japan). The energy dispersive spectroscopy (EDS) was taken at 20 kV and a vacuum of 104 Pa to investigate the arrangement of the crystalline phases. To reveal the morphology of the crystalline structure, an etching solution of hydrofluoric acid (HF acid, 5 wt.%) was used to remove the glassy phase in the sintered samples. 3. Results and discussion Fig. 2. XRD patterns of samples AS sintered at different temperatures.
3.1. Phase transformation Phase transformation of samples AS sintered in the temperature range 1460e1540 C was examined by XRD analysis (Fig. 2), and the relative content of the crystalline phases obtained at each temperature was summarized in Table 2. The dominant phases of samples AS sintered at 1460 C and 1480 C are SiC, mullite (Al4.75Si1.25O9.63) and corundum (Al2O3). The XRD patterns of mullite fit well with PDF 79e1454 (Al2O3:SiO2 ¼ 1.9:1 in mol%), indicating that mullite synthesized in the reducing atmosphere is Al2O3-rich mullite with 79.1 wt.% Al2O3. After the completion of the carbothermal reaction at 1500 C, SiC and Al2O3 are detected as final products, along with the complete decomposition of the intermediate mullite. As seen in Table 2, the SiC content of 54.3 wt.% in samples AS sintered at 1460 C is close to the calculated result in Table 1 with a little deviation of 1.97%. The SiC content increases with elevated temperature and approximates to the theoretical value of 64.83 wt.% when sintering at 1500 C. Compared with the finish temperature of 1550 C reported in Ref. [35] for the carbothermal reduction of andalusite-carbon system under Ar protection, the finish temperature in present study has been lowered by 50 C.
For exploring the effects of Y2O3 and SiC additions on the carbothermal reaction rate, the phase transformation in samples AS1 and AS2 sintered at 1440e1560 C was also studied by XRD (Fig. 3). Fig. 3a shows that the carbothermal reaction of samples AS1 finishes at 1440 C with the production of Al2O3 and 3CeSiC (b-SiC). Without the help of Y2O3, however, mullite in samples AS2 cannot decompose completely until 1560 C (Fig. 3b). Thus it is believed that the addition of Y2O3 promotes the carbothermal reduction of mullite and decreases the finish temperature to 1440 C, which shall be attributed to the positive effects of Y2O3 on the removal of oxygen and silicon atoms and the substitution of Si4þ by Al3þ. The ternary Y2O3eSiO2eAl2O3 system can form a metastable eutectic liquid with low viscosity at a lower temperature than SiO2eAl2O3 system (the eutectic melt of Y2O3 with SiO2eAl2O3 binary system forms at <1400 C) [36,37], thus inducing the solution of mullite into the liquid and increasing the diffusion rate of oxygen and silicon atoms simultaneously. The solution and diffusion processes allow the bounded silica in mullite to contact with the reducing agent (CO gas) directly. As a result, the reaction between silica and CO gas becomes more extensive at lower temperature, after which
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Table 2 The variations in the relative content of crystalline phases of samples AS with elevating temperature (wt %). Sintering temperature/ C
Silicon carbide
Mullite(Al4.75Si1.25O9.63)
Corundum
1460 1480 1500 1520 1540 Theoretical value
54.3 63.7 64.8 64.5 64.2 64.83
26.3 10.7 e e e e
19.5 25.6 35.2 35.5 35.8 35.17
The theoretical value is calculated by adding the amount of the in situ formed SiC and subtracting the amount of impurities.
Fig. 3. XRD patterns of reference samples heated at 1440e1560 C: (a) AS1 (with Y2O3 addition); (b) AS2 (without Y2O3 addition).
the liquid with low viscosity also assists Al3þ to move to substitute Si4þ. On the contrary, the added SiC exerts a negative effect on the carbothermal reaction, as the reaction in samples AS with SiC addition finished at higher temperature than samples AS1 without SiC addition (see Fig. 2 and Fig. 3a). The reason responsible for the higher finish temperature for samples AS shall be the smaller quantity of liquid phase. As known, mass transfer highly depends on the quantity of liquid phase when using the liquid phase sintering process [38]. On the other hand, it is discovered that mullite is reduced not only by carbon in the gaseous form of CO as suggested by Eq. (1), but also by the solid carbon penetrated from the buried carbon bed through reaction in Eq. (2) [35,37,39,40], which can be verified by the observation in Fig. 4: the buried carbon penetrates into sample AS2 from surface to interior with the unreacted materials in gray color (grey in the web version) when sintering at 1480 C, while the fractured surface of sample AS1 becomes uniform in black color (black in the web version) with the addition of Y2O3. This observation suggests the relationship between the carbothermal reaction rate and carbon penetration: the larger quantity of liquid phase, the faster penetration rate. Whereas the quantity of liquid phase in samples AS is decreased significantly by the addition of SiC, thus resulting in the slower carbon
Fig. 4. Image of fractured surface of samples AS1 and AS2 heated at 1480 C.
dissolution/diffusion processes and thus the slower reaction rate.
3Al2 O3 $2SiO2 ðsÞþ4SiO2 ðsÞ þ 18CðsÞ/3Al2 O3 ðsÞ þ 6SiCðsÞ þ 12COðgÞ (2) 3.2. Microstructural evolution SEM observations of fractured surface of samples AS sintered at 1460 C, 1500 C and 1540 C revealed the microstructural evolution with increasing temperature (Fig. 5). Morphology comparison from Fig. 5a, c and e with the same magnification indicates that the microstructure of samples AS becomes the densest when sintering at 1540 C, since the large voids observed in samples sintered at
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Fig. 5. SEM figures of fractured surface of samples AS sintered at different temperatures: (a)&(b) 1460 C; (c)&(d) 1500 C; (e)e(h) 1540 C. (e) and (f) refer to fractured surface of samples sintered at 1540 C after HF etching, while (g) and (h) refer to samples before HF etching. (d), (f) and (h) are the high magnifications of M zones in (c), (e) and (g), respectively.
1460 C and 1500 C vanish, which agrees well with testing results of water absorption and open porosity arranged below. Fig. 5f shows that the microstructure of samples AS sintered at 1540 C predominantly consists of SiC and Al2O3, and the SiC particles with the typical prismatic structure embed in the matrix constructed by Al2O3 crystals. This SEM observation serves as evidence that the
sintering mechanism for samples AS is the bonding of SiC by the combination of the in situ produced Al2O3 crystals, which is induced by the liquid phase sintering process at temperatures <1460 C. It is interesting to note in Fig. 5 that, many particles in blossomshape are located on the surfaces of SiC and Al2O3 crystals and those
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Fig. 6. EDS (Energy Dispersive Spectroscopy) patterns of three selected spots in Fig. 5f.
from the samples sintered at higher temperature have larger sizes (Fig. 5b, d and f). Differences in morphology observed between these blossom-shaped particles in the samples sintered at the higher temperature and the lower ones reveal that the blossomshaped particles grow with increasing temperature. EDS analysis results (Fig. 6 and Table 3) indicate that the elements of the Table 3 EDS analysis results of the blossom-shaped particles shown in Fig. 5f. Element
Spot 1
Spot 2
Spot 3
C O Al Si Y Fe Ca Total
28.50 0 0 71.50 0 0 0 100
62.16 16.28 4.61 13.96 2.99 0 0 100
50.52 25.71 9.10 9.44 4.16 0.46 0.58 100
blossom-shaped particles are dominated by Y, Al, Si, O and C, which are the basic elements of YeAleSieOeC glass [41]. In contrast to XRD results, since crystals of YeAleSieOeC system are undetected, these particles are probably glassy phases. Also the absence of the blossom-shaped particles in the morphology of samples without HF etching demonstrates that they are the residues of glassy phase after HF etching (Fig. 5g and h). On the other hand, the detection of carbide (C), calcium (Ca) and iron (Fe) in glassy phases confirms the carbon penetration process via liquid phase mass transfer. Nevertheless, the further growth of the blossom-shaped particles becomes unfavorable for physical properties. As seen in Fig. 7, the blossom-shaped particles grow abnormally to wrap up SiC and Al2O3 crystals when the sintering temperatures increase to 1560e1580 C, resulting in the weak bonding of the crystals. Moreover, the blossom-shaped particles as glassy phases show the intrinsic non-plastic characteristic with the formation of cracks on their surfaces (Fig. 7c and d). The weak bonding and the formed
Fig. 7. SEM figures of fractured surface of samples AS sintered at different temperatures: (a)&(b) samples sintered at 1560 C after HF etching; (c)&(d) samples sintered at 1580 C after HF etching. (d) and (f) are the high magnifications of M zones in (c) and (e), respectively.
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cracks cause the strength degradation of samples AS sintered at 1560e1580 C, which will be mentioned below. In order to reveal the effects of Y2O3 and SiC additions on the microstructural evolution, the SEM analysis of samples AS1 and AS2 was performed as well (Fig. 8). Fig. 8b and d reveals the presence of the blossom-shaped particles in samples AS1 (with the addition of Y2O3) and their absence in samples AS2 (without the addition of Y2O3). It is evident that Y2O3 plays an important role on the formation of the blossom-shaped particles. In addition, without the help of Y2O3, there are remains of mullite in samples AS2 sintered at 1480 C due to the slower carbothermal reaction rate (Fig. 8c and d). After sintering at a higher temperature as 1560 C, the bound silica in mullite is reduced to produce SiC whiskers (Fig. 8e and f), whose morphology and formation mechanism have been widely investigated. As reported in Refs. [42,43], SiC whiskers are yielded by the carbothermal reaction between silica and carbon via the following sequences:
SiO2 ðsÞ þ CðsÞ/SiOðgÞ þ COðgÞ
(3)
SiO2 ðsÞ þ COðgÞ/SiOðgÞþCO2 ðgÞ
(4)
SiOðgÞ þ 2CðsÞ/SiCðsÞ þ COðgÞ
(5)
Here silica reacts with solid carbon and CO gas to form gaseous silicon monoxide (SiO) through Eqs. (3) and (4). SiC whiskers are then produced via the reaction described in Eq. (5) between SiO and solid carbon. It is important to note that, the equilibrium condition at final step depends on the temperature and the partial pressure of SiO (PSiO) [43,44]. In case of formula AS, the amount of SiO2 is much smaller than that of formula AS2 due to the addition of SiC, thus leading to a small rate of SiO production and the low PSiO. Therefore, no SiC whisker can be observed in samples AS and SiC nanoparticles instead of whiskers shall be obtained. 3.3. Relevant variations in physical and thermophysical properties of Al2O3/SiC composite ceramics Fig. 9 shows the water absorption (Wa), open porosity (Pa), density (D) and bending strength of samples AS as the functions of sintering temperature. As seen in Fig. 9a, Wa and Pa decrease with the increase of sintering temperature until 1540 C and reach the lowest values of 11.0% and 26.3%, respectively, while bending strength and D increase with elevated temperature and achieve the highest values of 46.7 MPa and 2.37 g/cm3 when sintering at 1540 C (Fig. 9). The slight increase of Wa and Pa at 1500 C shall be attributed to the generation of gaseous SiO and CO during the
Fig. 8. SEM figures of fractured surface of reference samples: (a)&(b) samples AS1 sintered at 1480 C; (c)&(d) samples AS2 sintered at 1480 C; (e)&(f) samples AS2 sintered at 1560 C. (b), (d) and (f) are the high magnifications of M zones in (a), (c) and (e), respectively.
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Fig. 9. Water absorption(Wa), open porosity(Pa), density(D) and bending strength of samples AS as the functions of sintering temperature: (a) Wa and Pa; (b) bending strength and D.
carbothermal reaction (Eqs. (1) and (2)). According to XRD results, the carbothermal reaction becomes extensive at 1500 C, generating more gas to hinder densification process and thereby increasing the open porosity. Therefore, after the reaction finishing at 1520e1540 C, a smaller Wa, a smaller Pa, and a larger D can be obtained as the microstructure becomes denser. In contrast to SEM analysis, the drastic drop in bending strength above 1540 C is due to the weak bonding of SiC and Al2O3 and the non-plastic nature of the blossom-shaped particles. According to the results above, 1540 C is determined as the optimal sintering temperature for the obtainment of favorable physical properties, and samples AS sintered at this temperature were selected for the tests of thermophysical properties (Table 4). As seen, the heat capacity increases with increasing temperature from 0.68 J/(g$K) at room temperature to 1.16 J/(g$K) at 600 C, which indicates that the thermal storage capability of the composite ceramics will be enhanced at higher working temperatures. At room temperature, samples AS present the maximum thermal conductivity and thermal diffusivity values of 13.73 W/(m$K) and 8.44 mm2/s, respectively. Incorporation of the in situ formed Al2O3 to SiC ceramics clearly enhances the heat capacity, while the effective reduction of silica greatly contributes to improve the thermal conductivity. As seen in Table 5, it is testified that samples AS2 sintered at 1560 C show a higher heat capacity than samples AS (up to 1.18 J/(g$K) at 600 C), demonstrating that the in situ formed Al2O3 material is featured of high heat capacity. Due to the effective reduction of silica with low
thermal conductivity and the in situ synthesis of 3CeSiC with high thermal conductivity, the relatively high thermal conductivity of 6.97 W/(m$K) at 30.6 C is obtained for samples AS2, which is three times higher than that of the Al2O3-based ceramics (i.e.1.5e2.3 W/ (m$K)) reported in our previous studies [13,31,32]. Despite the high heat capacity of the in situ formed Al2O3, the heat capacity of samples AS is lower than the values calculated by Al2O3 and SiC weight fractions (see Table 6, the heat capacity data are from Ref. [45]). This discrepancy is probably related to the existence of impurities, such as the blossom-shaped particles. In terms of thermal conductivity, a higher value for samples AS shall be expected if considering the thermal conductivity value (70e80 W/(m$K)) for the porous SiC ceramics with a porosity of 30% [46]. However, the degradation for thermal conductivity cannot be avoided due to the incorporation of Al2O3, the presence of the impurities and interfacial defects (i.e. pores and grain boundaries). 3.4. Thermal shock resistance of Al2O3/SiC composite ceramics Generally, the thermal storage-retrieval process of CSP is in the temperature range 400e850 C using air as HTF to guarantee the power generation efficiency, thus the thermal storage materials are required to have a good thermal shock resistance to endure a longterm process. Table 7 presents the bending strength and the strength loss rate of samples after different cycles of thermal shock. There is little or no significant strength changes for any of samples AS after thermal shock, which suggests that the good thermal shock
Table 4 Thermophysical properties of samples AS sintered at 1540 C. Testing temperature/ C
Heat capacity /J$(g$K)1
Thermal diffusivity/mm2$s1
Thermal conductivity /W$(m$K)1
33.3 300 600
0.68 1.03 1.16
8.44 4.47 3.17
13.73 10.90 8.63
Table 5 Thermophysical properties of samples AS2 sintered at 1560 C. Testing temperature/ C
Heat capacity /J$(g$K)1
Thermal diffusivity/mm2$s1
Thermal conductivity /W$(m$K)1
30.6 300 600
0.81 1.08 1.18
4.21 2.19 1.55
6.97 4.81 3.72
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Table 6 Heat capacity values used in the calculations of Al2O3/SiC composites and the calculated results at different temperatures. Data were from Ref. [45]. Temperature ( C)
Cp SiC/J$(g$K)1
Cp Al2O3/J$(g$K)1
Calculated Cp (35.8%Al2O3-64.2%SiC)/J$(g$K)1
25 300 600
0.6628 1.0309 1.1606
0.7719 1.0906 1.1933
0.7018 1.0523 1.1723
Table 7 Bending strength and strength loss rate of samples AS after thermal shock. Cycle times of thermal shock
1
10
20
30
Bending strength of samples after thermal shock/MPa Strength loss rate/%
48.1 3.1
48.6 4.2
47.4 1.6
47.8 2.5
The negative note “-” refers to the increase of bending strength instead of decrease after thermal shock.
resistance is achieved for the Al2O3/SiC composite ceramics to meet the requirement of sustaining the thermal shock in the severe working condition. Fig. 10 shows the morphology of samples AS after different cycles of thermal shock. As observed in Fig. 10e and f, surfaces of
samples AS do not presents micro-crack deflection, grain growth and densification of microstructure even after 30 cycles of thermal shock, which are known as the enhancement mechanisms for bending strength. Thus, the good thermal shock resistance of samples AS shall depend on the relatively high thermal
Fig. 10. SEM figures of fractured surface of samples AS underwent different cycles of thermal shock: (a)&(b) 10 cycles; (c)&(d) 20 cycles; (e)&(f) 30 cycles.
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Fig. 11. EDS patterns of three selected spots in Fig. 10e and f.
conductivity, which is in favor of eliminating temperature gradient during thermal shock. According to the thermal-shock damage theory [47], which regards the equilibrium of thermo-elastic strain energy and fracture surface energy as the criterion of thermal shock resistance, when the temperature gradient causes a greater thermo-elastic strain energy than fracture surface energy, cracks will form and extend, finally resulting in the thermal shock failure of materials. Compared with the microstructure of unshocked ceramics in Fig. 5f, Fig. 10f shows that the blossom-shaped particles vanish and an amorphous film develops gradually on the surfaces of Al2O3 crystals. As observed in Fig. 10a and b, there are only a few residues of the blossom-shaped particles after 10 cycles of thermal shock; with increasing cycles of thermal shock, the blossom-shaped particles disappear completely and many spherical particles begin to form (Fig. 10c and d); these spherical particles then grow up to form the amorphous film after 30 cycles of thermal shock. EDS patterns (Fig. 11) show that the amorphous film is consistent with the blossom-shaped particles in terms of elemental composition, indicating that the generation of the amorphous film shall be ascribed to the collapse of the blossom-shaped particles. According to EDS semi-quantitative results listed in Table 8, carbon loss during Table 8 EDS analysis results of the selected spots shown in Fig. 10e and f.
the thermal shock at 1100 C may be the reason responsible for the collapse of the blossom-shaped particles since the carbon content of the amorphous film is much lower than that of the blossomshaped particles (Table 3). 3.5. Comparison of Al2O3/SiC composite ceramics with other reported thermal storage ceramic materials Above experimental results have been collected for the comparison with the potential thermal storage ceramics (Table 9) [48e51]. As seen in Table 9, the thermal conductivity of Al2O3/SiC composite ceramics shows notable lower value than that of pure SiC ceramic, although it approximates to that of hot-pressed Si3N4 and is much higher than the reported values of other thermal storage ceramics. Thus the major differentiating characteristics of the composite ceramics are the higher heat capacity, the lower sintering temperature, the simpler synthesis conditions, and the lower manufacturing cost in comparison to pure SiC ceramic. As compared with the oxide thermal storage ceramics, the major advantages of the composite ceramics for solar sensible thermal storage are the higher thermal conductivity, the higher bending strength and the better thermal shock resistance in a large working temperature range of room temperature to 1100 C. 4. Conclusions
Element
Spot 1
Spot 2
Spot 3
C O Al Si Y Total
24.77 0 0 75.22 0 100
9.20 39.40 23.50 16.21 11.66 100
22.29 29.63 18.74 25.33 4.00 100
A technique of removing silica in aluminosilicate precursor by carbothermal reduction method was used to in situ form Al2O3 to bond with SiC for synthesizing Al2O3/SiC composites at a relatively low temperature of 1540 C. The Al2O3/SiC composite ceramics were characterized by the high heat capacity, the high thermal conductivity and the good thermal shock resistance, which would
Table 9 The main properties of typical high temperature sensible thermal storage ceramics (room temperature) [48e51]. Name
Thermal conductivity (W/(m$K))
Heat capacity (J/(g$K))
Bending strength(MPa)
Density(g/cm3)
Al2O3/SiC composites SiC Si3N4(hot pressed) Corundum(Al2O3) Cordierite-andalusite Cordierite Mullite Al2O3eSiCeZrO2 composites Castable ceramic Zirconite ACW(asbestos -containing wastes) ceramics Silica bricks
13.73 70-80 14.0 2.2 1.12 1.97e2.32 5.2 1.5-2.3 1.0e1.5 3.2e3.5 2.1e1.4
0.68 at RT(room temperature) 0.66 at RT e 0.77 at RT e 0.35 at RT 0.76 at RT 0.5 at 200 C 0.86 at 350 C 0.71 0.8-1.03 at 200-1000 C 0.44 at RT
46.7 e 56.8 25 93 3e30 25 66.2 e 25 e
2.37 3.2 3.24 2.5e3.2 2.14 1.7 3.23 3.2e3.5 3.5 3.0 3.1
e
1.82
1.5
136
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make the composites suitable for solar sensible thermal storage. Several key conclusions can be drawn below. (1) The addition of Y2O3 increased the carbothermal reaction rate significantly to assist the composites to finish the carbothermal reaction at 1500 C. Y2O3 behaved as a catalyst to induce and favor the growth of glassy particles of YeAleSieOeC system at 1460e1540 C, of which the abnormal growth at 1560e1580 C became detrimental to the physical properties of the composites. (2) Al2O3 and SiC in situ formed from the carbothermal reduction of mullite could improve the heat capacity and the thermal conductivity of the composite ceramics, whilst the composite ceramics sintered at 1540 C exhibited the optimal physical and thermophysical properties, which were listed as follows: 46.7 MPa for bending strength, 11.0% for water absorption, 1.16 J/(g$K) for heat capacity, 13.73 W/(m$K) for thermal conductivity. The relatively high thermal conductivity endowed the composite ceramics with the good thermal shock resistance. (3) The advantages of the Al2O3/SiC composite ceramics for thermal storage applications were the lower sintering temperature, the simpler synthesis conditions, the lower manufacturing cost than SiC ceramic, and the higher thermal conductivity, the higher bending strength and the better thermal shock resistance compared with oxide ceramics. Acknowledgment Authors are very grateful to the financial support from “973 Program (2010CB227105)”, P. R. China. References [1] Dynamic simulation of thermal energy storage system of Badaling 1MW solar power tower plant, Renew. Energ 39 (2012) 455e462. [2] Y. Tian, C.Y. Zhao, A review of solar collectors and thermal energy storage in solar thermal applications, Appl. Energ 104 (2013) 538e553. [3] H. Singh, R.P. Saini, J.S. Saini, A review on packed bed solar energy storage systems, Renew. Sustain, Ener. Rev. 14 (2010) 1059e1069. , E. Oro , D. Boer, L.F. Cabeza, Embodied energy in thermal energy [4] L. Miro storage (TES) systems for high temperature applications, Appl. Energ 137 (2015) 793e799. [5] A.I. Fernandez, M. Martínez, M. Segarra, I. Martorell, L.F. Cabeza, Selection of materials with potential in sensible thermal energy storage, Sol. Energ. Mat. Sol. C 94 (2010) 1723e1729. [6] S. Kuravi, J. Trahan, D. Yogi Goswami, M.M. Rahman, E.K. Stefanakos, Thermal energy storage technologies and systems for concentrating solar power plants, Prog. Energ. Combust. 39 (2013) 285e319. , Eduard Oro , Dieter Boer, Luisa F. Cabeza, Embodied energy in [7] Laia Miro thermal energy storage (TES) systems for high temperature applications, Appl. Energ 137 (2015) 793e799. [8] S. Khare, M. Dell’Amico, C. Knight, S. McGarry, Selection of materials for high temperature sensible energy storage, Sol. Energ. Mat. Sol. C 115 (2013) 114e122. [9] G. Mebrahitom Asmelash, O. Mamat, F. Ahmad, A.K. Prasadarao, Thermal shock and fatigue behavior of pressureless sintered Al2O3-SiO2-ZrO2 composites, J. Advan. Ceram. 4 (2015) 190e198. [10] Y.X. Wang, M. Dong, Research development of the thermal shock resistance of thermal storage honeycomb ceramic, Chin. Ceram. 47 (2011) 1e6. [11] Y.J. Zhong, Q.G. Guo, S.Z. Li, J.L. Shi, L. Liu, Heat transfer enhancement of paraffin wax using carbon foam for thermal energy storage, Sol. Energ. Mat. Sol. C 94 (2010) 1011e1014. [12] G. Zanganeh, A. Pedretti, A. Haselbacher, A. Steinfeld, Design of packed bed thermal energy storage systems for high-temperature industrial process heat, Appl. Eng. 137 (2015) 812e822. [13] X.H. Xu, F. Zhao, J.F. Wu, J. Li, C.G. Li, Al2O3/SiC composite ceramics used in solar thermal power, J. Wuhan. Univ. Technol. 31 (13) (2009) 8e11. [14] R.X. Tan, Z.L. Ke, Y.C. Liu, S.A. Ze, H. Wang, P.A. Liu, J. Song, Porous ceramic materials used for heat storage and accumulation, J. Ceram. 2 (27) (2006) 217e226. [15] S.Q. Ding, Y.P. Zeng, D.L. Jiang, Thermal shock behavior of mullite-bonded porous silicon carbide ceramics with yttria addition, J. Phys. D. Appl. Phys. 40 (2007) 2138e2142.
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Contents lists available at ScienceDirect
Journal of Alloys and Compounds journal homepage: http://www.elsevier.com/locate/jalcom
Synthesis and characterization of Al2O3/SiC composite ceramics via carbothermal reduction of aluminosilicate precursor for solar sensible thermal storage Xiaohong Xu, Xinbin Lao*, Jianfeng Wu, Yaxiang Zhang, Xiaoyang Xu, Kun Li State Key Laboratory of Silicate Materials for Architectures, Wuhan University of Technology, Wuhan 430070, PR China
a r t i c l e i n f o
a b s t r a c t
Article history: Received 25 August 2015 Received in revised form 2 December 2015 Accepted 7 December 2015 Available online 12 December 2015
Al2O3 material with high heat capacity and high thermal conductivity was in situ formed to bond with SiC for synthesizing Al2O3/SiC composite ceramics by the removal of silica in aluminosilicate precursor using carbothermal reduction method. Phase transformation, microstructural evolution, and the relevant variations in properties were studied in the temperature range 1460e1580 C. Results indicated that the addition of Y2O3 improved the carbothermal reaction rate significantly and decreased the finish temperature of carbothermal reaction to 1500 C. Incorporation of the in situ formed Al2O3 could lower the sintering temperature of SiC ceramic without decreasing the heat capacity. By sintering at the optimal temperature as 1540 C, Al2O3/SiC composite ceramics with the heat capacity of 1.16 J/(g$K) and the high thermal conductivity of 13.73 W/(m$K) were obtained. The relatively high thermal conductivity endowed the composites with a good thermal shock resistance. This study was intended to identify the reaction conditions for obtaining sensible thermal storage materials with favorable morphology and properties by using a simple in situ synthesis method. © 2015 Elsevier B.V. All rights reserved.
Keywords: Al2O3/SiC composite ceramics Solar sensible thermal storage Carbothermal reduction Heat capacity Thermal conductivity
1. Introduction CO2-induced global warming has become a pressing issue, and needs to be tackled urgently. Efficient utilization of solar energy is thereby proposed and being considered increasingly as a way to achieve a sustainable and clean energy supply for human beings [1,2]. Solar thermal power generation is one of the prevalent power generation technologies using solar energy, and the conversion of solar energy into thermal energy is the easiest and the most extensively accepted method [3]. The main problem of solar thermal power generation is the mismatch between the solar energy and the energy demand, since solar energy is an intermittent energy resource. To remove the fluctuations caused by the intermittent nature of solar energy, a thermal storage system composed of thermal storage materials is required to be attached with solar collectors to store energy and ensure the continuous power supply [4,5]. The thermal storage materials are classified into three main categories according to different storage technologies: sensible heat storage (SHS), latent heat storage, thermo-chemical storage.
* Corresponding author. E-mail address: [email protected] (X. Lao). http://dx.doi.org/10.1016/j.jallcom.2015.12.038 0925-8388/© 2015 Elsevier B.V. All rights reserved.
Among these thermal storage technologies, sensible heat storage is the most simple and inexpensive way for thermal storage system with technological and economical superiority, although there are few advantages of latent storage and thermo-chemical storage over sensible heat storage [6,7]. Excellent properties of the thermal storage materials are the key factors to guarantee the operation of solar sensible thermal storage system. A high thermal storage capacity (i.e. sensible heat capacity) is essential to reduce the volume of sensible thermal storage system and increase the thermal storage efficiency, whilst a good thermal shock resistance allows the thermal storage materials to avoid mechanical degradation after thermal cycles [2]. Numerous oxide ceramics, such as alumina, zirconia, mullite, cordierite, were discovered to be suitable for SHS due to the high thermal storage capacity, the good thermo-chemical stability and the good corrosion resistance against hot heat transfer fluid (HTF, usually air) [8]. Owing to the poor thermal conductivity, however, oxide ceramics were susceptible to thermal shock, and the drastic degradation in mechanical properties after the thermal shock had limited their wide applications at high temperature (850 C), especially the thermal storage application [9,10]. Also the poor thermal conductivity of the oxide ceramics led to the poor heat transfer capability,
X. Xu et al. / Journal of Alloys and Compounds 662 (2016) 126e137
which might slow down thermal storage-retrieval process of the thermal storage system [11,12]. SiC ceramics consequently attracted much attention as candidate materials for solar sensible thermal storage due to the high thermal conductivity as well as the high thermal storage capacity [13,14]. The high sintering temperature (2100 C) of covalent compound SiC [15,16] and the high cost of SiC starting materials [17], however, rendered the difficulties of producing SiC ceramics with low manufacturing cost. Due to the high sintering temperature, the complex sintering processes of SiC ceramics, such as the hot-pressing sintering [18], the spark plasma sintering [19,20], and the vacuum sintering [21], etc, were generally employed for the production of SiC ceramics, which imposed strict requirement on sintering equipment and made the fabrication of products with complex shapes extremely difficult [22], thus hampering the comprehensive utilization of SiC ceramics in solar sensible thermal storage. As known, pressureless liquid phase sintering of SiC by introducing oxide materials as sintering additives did not suffer from these shortcomings, and the introduced oxide materials (e.g. Al2O3, silica, and mullite, etc.) not only promoted the sintering of SiC, but also substituted partial SiC to synthesize SiC composite ceramics with lower cost [23e26]. Among the commonly-used oxide materials, Al2O3 possesses optimal thermophysical properties. For example, in the temperature range 25e600 C, the heat capacity of Al2O3 is 0.77e1.19 J/(g$K), while those of silica and mullite are 0.44e1.13 J/(g$K), 0.76e1.18 J/(g$K), respectively [27,28]; the thermal conductivity of Al2O3 is 17.5 W/ (m$K) at room temperature, which exceeds the thermal conductivity levels of silica and mullite (i.e. 1.4 W/(m$K) for silica, 6.1 W/ (m$K) for mullite) [29e31]. In our previous works [13,31,32], Al2O3eSiC composite ceramics with the high thermal storage capacity, such as Al2O3eSiC and Al2O3eSiCeZrO2 composites, were obtained by using aluminosilicate minerals (e.g. kaolin, andalusite, talc, etc.) as additives to promote the pressureless liquid phase sintering process, whilst the properties were characterized for elevating the suitability for solar sensible thermal storage. The composite ceramics suffered from the poor thermal shock resistance as the bending strength decreased significantly after 30 cycles of thermal shock (instantaneous temperature up to 1000 C), although the high thermal storage capacity and the low sintering temperature were accomplished. The poor thermal shock resistance of the composites was attributed to the poor thermal conductivity of 1.5e2.3 W/(m$K), which mainly resulted from the introduction of silica from the aluminosilicate additives. Hence there was an urgent need to remove the negative effect of silica on the thermal conductivity when using aluminosilicate minerals of low cost as raw materials for fabricating thermal storage materials. It is noteworthy that, as reported by R. Naghizadeh and F. Golestani-fard [33], silica in the aluminosilicates could be reduced in the form of gaseous silicon monoxide (SiO) by carbothermal reduction, whose principle was described in Eq. (1). This method is thereby considered to be an effective way to cope with the thermal conductivity degradation caused by the introduced silica.
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the composite ceramics were investigated, whilst the relevant variations in properties were studied in detail. The Al2O3/SiC composite ceramics were characterized by the high thermal storage capacity, the high thermal conductivity as well as the good thermal shock resistance. The present work was aimed at studying the reaction conditions for obtaining Al2O3/SiC composite thermal storage ceramics with favorable properties by a less complex and lowcost in situ synthesis method. 2. Material and methods 2.1. Starting materials Starting materials for experiments were SiC powder with two particle sizes of 106 mm and 58 mm (Dan Jiangkou Hongyuan Silicon Carbide Co., Ltd, Hubei, China), andalusite (under 250-mesh, Xinjiang Baoan New Energy-Mineral Co.,Ltd, Xinjiang, China), Al powder (under 250-mesh, Sinopharm Chemical Reagent Co.,Ltd, Shanghai, China) and Y2O3 (chemical pure, Sinopharm Chemical Reagent Co.,Ltd, Shanghai, China). Based on our previous study [31], the weight ratio of SiC to Al2O3 in the final products were taken as 65:35, respectively (i.e. 35 wt.% SiC was substituted by Al2O3). According to Eq. (1), the proportions of the starting materials for formula AS were calculated to be: 63.5 wt.% SiC, 22.7 wt.% andalusite, 13.8 wt.% Al. 3 wt.% Y2O3 was added as the catalyst for the decomposition of mullite as well as the sintering additive for densification, since the finish temperature for carbothermal reaction reported in Ref. [33] was as high as 1650 C. Al powder was added to react with the excess silica of andalusite to produce the intermediate mullite. The chemical compositions of starting materials and formula AS were listed in Table 1, in which the compositions of formula AS were calculated including Al2O3 derived from the transformation of Al. The other two formulae, i.e. AS1 and AS2, were designed to study the effects of Y2O3 and the added SiC on the carbothermal reaction rate and the microstructural evolution by performing XRD and SEM analysis of as-sintered samples. Both formulae excluded the addition of SiC, and the ratio of andalusite:Al was consistent with that of formula AS. Formulae AS1 and AS2 were differentiated by 8 wt.% Y2O3 addition and no Y2O3 addition, respectively. 2.2. Processing The starting materials were mixed thoroughly with 5 wt.% PVA binder addition by ball milling, and then pressed by a uniaxial pressure of 20 MPa into rectangular samples of 6.8 mm 6.8 mm 38 mm and cylindrical samples of F10 2 mm. The formed powder compacts were dried at 100 C for 24 h in an electric oven and pressureless sintered by the carbon-buried sintering method (Fig. 1) from 1440 C to 1580 C in a seggar full of carbon at molybdenum disiliciade furnace with a heating rate of 5 C/min and a holding time of 3 h at the maximum temperatures. 2.3. Characterization
3Al2 O3 $2SiO2 ðsÞ þ 2COðgÞ/3Al2 O3 ðsÞ þ 2SiOðgÞþ2CO2 ðgÞ (1) In this study, the carbothermal reduction method was employed to remove the negative effects of silica on the thermal conductivity as well as the heat capacity of Al2O3/SiC composite ceramics, which were fabricated using andalusite, an abundant aluminosilicate mineral with high alumina content, thus avoiding the conventional costly, energy-consuming and complicated sintering processes of SiC ceramics. The main factors affecting the carbothermal reaction rate, the phase transformation and the microstructural evolution of
The water absorption (Wa), the open porosity (Pa) and the bulk density (D) were measured by with the rectangular samples by AUY120 electronic analytical balance (Japanese Shimazu) through static weighing method. The bending strength was measured with the rectangular samples by Computer Control Electronic Universal Test Machine (Shenzhen Reger Instrument Co., Ltd., Guangdong, China). The heat capacity, the thermal diffusivity as well as the thermal conductivity were measured using the cylindrical samples by TC-7000H laser flash thermal constant analyzer(ULVAC SINKURIKO. Inc., Yokohama, Japan). The thermal shock resistance of the
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Table 1 Chemical compositions of starting materials and formula AS (wt%). Starting material
SiC
SiO2
Al2O3
Fe2O3
TiO2
CaO
MgO
K2O
Na2O
I.L
Total
Silicon carbide Andalusite Formula AS
99.19 0 56.27
0 43.25 8.77
0.17 50.87 33.7
0.19 1.11 0.33
0 0.16 0.03
0.09 1.28 0.31
0 0.14 0.03
0 0.74 0.15
0 0.53 0.11
0 1.49 0.3
99.63 99.57 100.0
Fig. 1. Systematic representation of pressureless carbon-buried sintering.
samples was tested under a thermal cycle described as follows: the samples were placed in a furnace and heated to the preset of temperature (1100 C), soaked for 20 min and then were taken out with cooling by wind (at a speed of 4 m/s) to room temperature (25 C). The bending strength loss rate of the samples after thermal shock tests represented the thermal shock resistance. Reported values were available from the average of at least 5 samples. X-ray diffraction (XRD) patterns were collected using a D/MAXIIIAdiffractometer (Rigaku Corporation, Japan) equipped with Cu Ka (l ¼ 1.54 Å) radiation. The relative content of different phases was calculated by RIR (Reference Intensity Ratio) method semiquantitatively [34]. The morphology was observed by SEM (JEOL Ltd., Japan). The energy dispersive spectroscopy (EDS) was taken at 20 kV and a vacuum of 104 Pa to investigate the arrangement of the crystalline phases. To reveal the morphology of the crystalline structure, an etching solution of hydrofluoric acid (HF acid, 5 wt.%) was used to remove the glassy phase in the sintered samples. 3. Results and discussion Fig. 2. XRD patterns of samples AS sintered at different temperatures.
3.1. Phase transformation Phase transformation of samples AS sintered in the temperature range 1460e1540 C was examined by XRD analysis (Fig. 2), and the relative content of the crystalline phases obtained at each temperature was summarized in Table 2. The dominant phases of samples AS sintered at 1460 C and 1480 C are SiC, mullite (Al4.75Si1.25O9.63) and corundum (Al2O3). The XRD patterns of mullite fit well with PDF 79e1454 (Al2O3:SiO2 ¼ 1.9:1 in mol%), indicating that mullite synthesized in the reducing atmosphere is Al2O3-rich mullite with 79.1 wt.% Al2O3. After the completion of the carbothermal reaction at 1500 C, SiC and Al2O3 are detected as final products, along with the complete decomposition of the intermediate mullite. As seen in Table 2, the SiC content of 54.3 wt.% in samples AS sintered at 1460 C is close to the calculated result in Table 1 with a little deviation of 1.97%. The SiC content increases with elevated temperature and approximates to the theoretical value of 64.83 wt.% when sintering at 1500 C. Compared with the finish temperature of 1550 C reported in Ref. [35] for the carbothermal reduction of andalusite-carbon system under Ar protection, the finish temperature in present study has been lowered by 50 C.
For exploring the effects of Y2O3 and SiC additions on the carbothermal reaction rate, the phase transformation in samples AS1 and AS2 sintered at 1440e1560 C was also studied by XRD (Fig. 3). Fig. 3a shows that the carbothermal reaction of samples AS1 finishes at 1440 C with the production of Al2O3 and 3CeSiC (b-SiC). Without the help of Y2O3, however, mullite in samples AS2 cannot decompose completely until 1560 C (Fig. 3b). Thus it is believed that the addition of Y2O3 promotes the carbothermal reduction of mullite and decreases the finish temperature to 1440 C, which shall be attributed to the positive effects of Y2O3 on the removal of oxygen and silicon atoms and the substitution of Si4þ by Al3þ. The ternary Y2O3eSiO2eAl2O3 system can form a metastable eutectic liquid with low viscosity at a lower temperature than SiO2eAl2O3 system (the eutectic melt of Y2O3 with SiO2eAl2O3 binary system forms at <1400 C) [36,37], thus inducing the solution of mullite into the liquid and increasing the diffusion rate of oxygen and silicon atoms simultaneously. The solution and diffusion processes allow the bounded silica in mullite to contact with the reducing agent (CO gas) directly. As a result, the reaction between silica and CO gas becomes more extensive at lower temperature, after which
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Table 2 The variations in the relative content of crystalline phases of samples AS with elevating temperature (wt %). Sintering temperature/ C
Silicon carbide
Mullite(Al4.75Si1.25O9.63)
Corundum
1460 1480 1500 1520 1540 Theoretical value
54.3 63.7 64.8 64.5 64.2 64.83
26.3 10.7 e e e e
19.5 25.6 35.2 35.5 35.8 35.17
The theoretical value is calculated by adding the amount of the in situ formed SiC and subtracting the amount of impurities.
Fig. 3. XRD patterns of reference samples heated at 1440e1560 C: (a) AS1 (with Y2O3 addition); (b) AS2 (without Y2O3 addition).
the liquid with low viscosity also assists Al3þ to move to substitute Si4þ. On the contrary, the added SiC exerts a negative effect on the carbothermal reaction, as the reaction in samples AS with SiC addition finished at higher temperature than samples AS1 without SiC addition (see Fig. 2 and Fig. 3a). The reason responsible for the higher finish temperature for samples AS shall be the smaller quantity of liquid phase. As known, mass transfer highly depends on the quantity of liquid phase when using the liquid phase sintering process [38]. On the other hand, it is discovered that mullite is reduced not only by carbon in the gaseous form of CO as suggested by Eq. (1), but also by the solid carbon penetrated from the buried carbon bed through reaction in Eq. (2) [35,37,39,40], which can be verified by the observation in Fig. 4: the buried carbon penetrates into sample AS2 from surface to interior with the unreacted materials in gray color (grey in the web version) when sintering at 1480 C, while the fractured surface of sample AS1 becomes uniform in black color (black in the web version) with the addition of Y2O3. This observation suggests the relationship between the carbothermal reaction rate and carbon penetration: the larger quantity of liquid phase, the faster penetration rate. Whereas the quantity of liquid phase in samples AS is decreased significantly by the addition of SiC, thus resulting in the slower carbon
Fig. 4. Image of fractured surface of samples AS1 and AS2 heated at 1480 C.
dissolution/diffusion processes and thus the slower reaction rate.
3Al2 O3 $2SiO2 ðsÞþ4SiO2 ðsÞ þ 18CðsÞ/3Al2 O3 ðsÞ þ 6SiCðsÞ þ 12COðgÞ (2) 3.2. Microstructural evolution SEM observations of fractured surface of samples AS sintered at 1460 C, 1500 C and 1540 C revealed the microstructural evolution with increasing temperature (Fig. 5). Morphology comparison from Fig. 5a, c and e with the same magnification indicates that the microstructure of samples AS becomes the densest when sintering at 1540 C, since the large voids observed in samples sintered at
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Fig. 5. SEM figures of fractured surface of samples AS sintered at different temperatures: (a)&(b) 1460 C; (c)&(d) 1500 C; (e)e(h) 1540 C. (e) and (f) refer to fractured surface of samples sintered at 1540 C after HF etching, while (g) and (h) refer to samples before HF etching. (d), (f) and (h) are the high magnifications of M zones in (c), (e) and (g), respectively.
1460 C and 1500 C vanish, which agrees well with testing results of water absorption and open porosity arranged below. Fig. 5f shows that the microstructure of samples AS sintered at 1540 C predominantly consists of SiC and Al2O3, and the SiC particles with the typical prismatic structure embed in the matrix constructed by Al2O3 crystals. This SEM observation serves as evidence that the
sintering mechanism for samples AS is the bonding of SiC by the combination of the in situ produced Al2O3 crystals, which is induced by the liquid phase sintering process at temperatures <1460 C. It is interesting to note in Fig. 5 that, many particles in blossomshape are located on the surfaces of SiC and Al2O3 crystals and those
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Fig. 6. EDS (Energy Dispersive Spectroscopy) patterns of three selected spots in Fig. 5f.
from the samples sintered at higher temperature have larger sizes (Fig. 5b, d and f). Differences in morphology observed between these blossom-shaped particles in the samples sintered at the higher temperature and the lower ones reveal that the blossomshaped particles grow with increasing temperature. EDS analysis results (Fig. 6 and Table 3) indicate that the elements of the Table 3 EDS analysis results of the blossom-shaped particles shown in Fig. 5f. Element
Spot 1
Spot 2
Spot 3
C O Al Si Y Fe Ca Total
28.50 0 0 71.50 0 0 0 100
62.16 16.28 4.61 13.96 2.99 0 0 100
50.52 25.71 9.10 9.44 4.16 0.46 0.58 100
blossom-shaped particles are dominated by Y, Al, Si, O and C, which are the basic elements of YeAleSieOeC glass [41]. In contrast to XRD results, since crystals of YeAleSieOeC system are undetected, these particles are probably glassy phases. Also the absence of the blossom-shaped particles in the morphology of samples without HF etching demonstrates that they are the residues of glassy phase after HF etching (Fig. 5g and h). On the other hand, the detection of carbide (C), calcium (Ca) and iron (Fe) in glassy phases confirms the carbon penetration process via liquid phase mass transfer. Nevertheless, the further growth of the blossom-shaped particles becomes unfavorable for physical properties. As seen in Fig. 7, the blossom-shaped particles grow abnormally to wrap up SiC and Al2O3 crystals when the sintering temperatures increase to 1560e1580 C, resulting in the weak bonding of the crystals. Moreover, the blossom-shaped particles as glassy phases show the intrinsic non-plastic characteristic with the formation of cracks on their surfaces (Fig. 7c and d). The weak bonding and the formed
Fig. 7. SEM figures of fractured surface of samples AS sintered at different temperatures: (a)&(b) samples sintered at 1560 C after HF etching; (c)&(d) samples sintered at 1580 C after HF etching. (d) and (f) are the high magnifications of M zones in (c) and (e), respectively.
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cracks cause the strength degradation of samples AS sintered at 1560e1580 C, which will be mentioned below. In order to reveal the effects of Y2O3 and SiC additions on the microstructural evolution, the SEM analysis of samples AS1 and AS2 was performed as well (Fig. 8). Fig. 8b and d reveals the presence of the blossom-shaped particles in samples AS1 (with the addition of Y2O3) and their absence in samples AS2 (without the addition of Y2O3). It is evident that Y2O3 plays an important role on the formation of the blossom-shaped particles. In addition, without the help of Y2O3, there are remains of mullite in samples AS2 sintered at 1480 C due to the slower carbothermal reaction rate (Fig. 8c and d). After sintering at a higher temperature as 1560 C, the bound silica in mullite is reduced to produce SiC whiskers (Fig. 8e and f), whose morphology and formation mechanism have been widely investigated. As reported in Refs. [42,43], SiC whiskers are yielded by the carbothermal reaction between silica and carbon via the following sequences:
SiO2 ðsÞ þ CðsÞ/SiOðgÞ þ COðgÞ
(3)
SiO2 ðsÞ þ COðgÞ/SiOðgÞþCO2 ðgÞ
(4)
SiOðgÞ þ 2CðsÞ/SiCðsÞ þ COðgÞ
(5)
Here silica reacts with solid carbon and CO gas to form gaseous silicon monoxide (SiO) through Eqs. (3) and (4). SiC whiskers are then produced via the reaction described in Eq. (5) between SiO and solid carbon. It is important to note that, the equilibrium condition at final step depends on the temperature and the partial pressure of SiO (PSiO) [43,44]. In case of formula AS, the amount of SiO2 is much smaller than that of formula AS2 due to the addition of SiC, thus leading to a small rate of SiO production and the low PSiO. Therefore, no SiC whisker can be observed in samples AS and SiC nanoparticles instead of whiskers shall be obtained. 3.3. Relevant variations in physical and thermophysical properties of Al2O3/SiC composite ceramics Fig. 9 shows the water absorption (Wa), open porosity (Pa), density (D) and bending strength of samples AS as the functions of sintering temperature. As seen in Fig. 9a, Wa and Pa decrease with the increase of sintering temperature until 1540 C and reach the lowest values of 11.0% and 26.3%, respectively, while bending strength and D increase with elevated temperature and achieve the highest values of 46.7 MPa and 2.37 g/cm3 when sintering at 1540 C (Fig. 9). The slight increase of Wa and Pa at 1500 C shall be attributed to the generation of gaseous SiO and CO during the
Fig. 8. SEM figures of fractured surface of reference samples: (a)&(b) samples AS1 sintered at 1480 C; (c)&(d) samples AS2 sintered at 1480 C; (e)&(f) samples AS2 sintered at 1560 C. (b), (d) and (f) are the high magnifications of M zones in (a), (c) and (e), respectively.
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Fig. 9. Water absorption(Wa), open porosity(Pa), density(D) and bending strength of samples AS as the functions of sintering temperature: (a) Wa and Pa; (b) bending strength and D.
carbothermal reaction (Eqs. (1) and (2)). According to XRD results, the carbothermal reaction becomes extensive at 1500 C, generating more gas to hinder densification process and thereby increasing the open porosity. Therefore, after the reaction finishing at 1520e1540 C, a smaller Wa, a smaller Pa, and a larger D can be obtained as the microstructure becomes denser. In contrast to SEM analysis, the drastic drop in bending strength above 1540 C is due to the weak bonding of SiC and Al2O3 and the non-plastic nature of the blossom-shaped particles. According to the results above, 1540 C is determined as the optimal sintering temperature for the obtainment of favorable physical properties, and samples AS sintered at this temperature were selected for the tests of thermophysical properties (Table 4). As seen, the heat capacity increases with increasing temperature from 0.68 J/(g$K) at room temperature to 1.16 J/(g$K) at 600 C, which indicates that the thermal storage capability of the composite ceramics will be enhanced at higher working temperatures. At room temperature, samples AS present the maximum thermal conductivity and thermal diffusivity values of 13.73 W/(m$K) and 8.44 mm2/s, respectively. Incorporation of the in situ formed Al2O3 to SiC ceramics clearly enhances the heat capacity, while the effective reduction of silica greatly contributes to improve the thermal conductivity. As seen in Table 5, it is testified that samples AS2 sintered at 1560 C show a higher heat capacity than samples AS (up to 1.18 J/(g$K) at 600 C), demonstrating that the in situ formed Al2O3 material is featured of high heat capacity. Due to the effective reduction of silica with low
thermal conductivity and the in situ synthesis of 3CeSiC with high thermal conductivity, the relatively high thermal conductivity of 6.97 W/(m$K) at 30.6 C is obtained for samples AS2, which is three times higher than that of the Al2O3-based ceramics (i.e.1.5e2.3 W/ (m$K)) reported in our previous studies [13,31,32]. Despite the high heat capacity of the in situ formed Al2O3, the heat capacity of samples AS is lower than the values calculated by Al2O3 and SiC weight fractions (see Table 6, the heat capacity data are from Ref. [45]). This discrepancy is probably related to the existence of impurities, such as the blossom-shaped particles. In terms of thermal conductivity, a higher value for samples AS shall be expected if considering the thermal conductivity value (70e80 W/(m$K)) for the porous SiC ceramics with a porosity of 30% [46]. However, the degradation for thermal conductivity cannot be avoided due to the incorporation of Al2O3, the presence of the impurities and interfacial defects (i.e. pores and grain boundaries). 3.4. Thermal shock resistance of Al2O3/SiC composite ceramics Generally, the thermal storage-retrieval process of CSP is in the temperature range 400e850 C using air as HTF to guarantee the power generation efficiency, thus the thermal storage materials are required to have a good thermal shock resistance to endure a longterm process. Table 7 presents the bending strength and the strength loss rate of samples after different cycles of thermal shock. There is little or no significant strength changes for any of samples AS after thermal shock, which suggests that the good thermal shock
Table 4 Thermophysical properties of samples AS sintered at 1540 C. Testing temperature/ C
Heat capacity /J$(g$K)1
Thermal diffusivity/mm2$s1
Thermal conductivity /W$(m$K)1
33.3 300 600
0.68 1.03 1.16
8.44 4.47 3.17
13.73 10.90 8.63
Table 5 Thermophysical properties of samples AS2 sintered at 1560 C. Testing temperature/ C
Heat capacity /J$(g$K)1
Thermal diffusivity/mm2$s1
Thermal conductivity /W$(m$K)1
30.6 300 600
0.81 1.08 1.18
4.21 2.19 1.55
6.97 4.81 3.72
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Table 6 Heat capacity values used in the calculations of Al2O3/SiC composites and the calculated results at different temperatures. Data were from Ref. [45]. Temperature ( C)
Cp SiC/J$(g$K)1
Cp Al2O3/J$(g$K)1
Calculated Cp (35.8%Al2O3-64.2%SiC)/J$(g$K)1
25 300 600
0.6628 1.0309 1.1606
0.7719 1.0906 1.1933
0.7018 1.0523 1.1723
Table 7 Bending strength and strength loss rate of samples AS after thermal shock. Cycle times of thermal shock
1
10
20
30
Bending strength of samples after thermal shock/MPa Strength loss rate/%
48.1 3.1
48.6 4.2
47.4 1.6
47.8 2.5
The negative note “-” refers to the increase of bending strength instead of decrease after thermal shock.
resistance is achieved for the Al2O3/SiC composite ceramics to meet the requirement of sustaining the thermal shock in the severe working condition. Fig. 10 shows the morphology of samples AS after different cycles of thermal shock. As observed in Fig. 10e and f, surfaces of
samples AS do not presents micro-crack deflection, grain growth and densification of microstructure even after 30 cycles of thermal shock, which are known as the enhancement mechanisms for bending strength. Thus, the good thermal shock resistance of samples AS shall depend on the relatively high thermal
Fig. 10. SEM figures of fractured surface of samples AS underwent different cycles of thermal shock: (a)&(b) 10 cycles; (c)&(d) 20 cycles; (e)&(f) 30 cycles.
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Fig. 11. EDS patterns of three selected spots in Fig. 10e and f.
conductivity, which is in favor of eliminating temperature gradient during thermal shock. According to the thermal-shock damage theory [47], which regards the equilibrium of thermo-elastic strain energy and fracture surface energy as the criterion of thermal shock resistance, when the temperature gradient causes a greater thermo-elastic strain energy than fracture surface energy, cracks will form and extend, finally resulting in the thermal shock failure of materials. Compared with the microstructure of unshocked ceramics in Fig. 5f, Fig. 10f shows that the blossom-shaped particles vanish and an amorphous film develops gradually on the surfaces of Al2O3 crystals. As observed in Fig. 10a and b, there are only a few residues of the blossom-shaped particles after 10 cycles of thermal shock; with increasing cycles of thermal shock, the blossom-shaped particles disappear completely and many spherical particles begin to form (Fig. 10c and d); these spherical particles then grow up to form the amorphous film after 30 cycles of thermal shock. EDS patterns (Fig. 11) show that the amorphous film is consistent with the blossom-shaped particles in terms of elemental composition, indicating that the generation of the amorphous film shall be ascribed to the collapse of the blossom-shaped particles. According to EDS semi-quantitative results listed in Table 8, carbon loss during Table 8 EDS analysis results of the selected spots shown in Fig. 10e and f.
the thermal shock at 1100 C may be the reason responsible for the collapse of the blossom-shaped particles since the carbon content of the amorphous film is much lower than that of the blossomshaped particles (Table 3). 3.5. Comparison of Al2O3/SiC composite ceramics with other reported thermal storage ceramic materials Above experimental results have been collected for the comparison with the potential thermal storage ceramics (Table 9) [48e51]. As seen in Table 9, the thermal conductivity of Al2O3/SiC composite ceramics shows notable lower value than that of pure SiC ceramic, although it approximates to that of hot-pressed Si3N4 and is much higher than the reported values of other thermal storage ceramics. Thus the major differentiating characteristics of the composite ceramics are the higher heat capacity, the lower sintering temperature, the simpler synthesis conditions, and the lower manufacturing cost in comparison to pure SiC ceramic. As compared with the oxide thermal storage ceramics, the major advantages of the composite ceramics for solar sensible thermal storage are the higher thermal conductivity, the higher bending strength and the better thermal shock resistance in a large working temperature range of room temperature to 1100 C. 4. Conclusions
Element
Spot 1
Spot 2
Spot 3
C O Al Si Y Total
24.77 0 0 75.22 0 100
9.20 39.40 23.50 16.21 11.66 100
22.29 29.63 18.74 25.33 4.00 100
A technique of removing silica in aluminosilicate precursor by carbothermal reduction method was used to in situ form Al2O3 to bond with SiC for synthesizing Al2O3/SiC composites at a relatively low temperature of 1540 C. The Al2O3/SiC composite ceramics were characterized by the high heat capacity, the high thermal conductivity and the good thermal shock resistance, which would
Table 9 The main properties of typical high temperature sensible thermal storage ceramics (room temperature) [48e51]. Name
Thermal conductivity (W/(m$K))
Heat capacity (J/(g$K))
Bending strength(MPa)
Density(g/cm3)
Al2O3/SiC composites SiC Si3N4(hot pressed) Corundum(Al2O3) Cordierite-andalusite Cordierite Mullite Al2O3eSiCeZrO2 composites Castable ceramic Zirconite ACW(asbestos -containing wastes) ceramics Silica bricks
13.73 70-80 14.0 2.2 1.12 1.97e2.32 5.2 1.5-2.3 1.0e1.5 3.2e3.5 2.1e1.4
0.68 at RT(room temperature) 0.66 at RT e 0.77 at RT e 0.35 at RT 0.76 at RT 0.5 at 200 C 0.86 at 350 C 0.71 0.8-1.03 at 200-1000 C 0.44 at RT
46.7 e 56.8 25 93 3e30 25 66.2 e 25 e
2.37 3.2 3.24 2.5e3.2 2.14 1.7 3.23 3.2e3.5 3.5 3.0 3.1
e
1.82
1.5
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make the composites suitable for solar sensible thermal storage. Several key conclusions can be drawn below. (1) The addition of Y2O3 increased the carbothermal reaction rate significantly to assist the composites to finish the carbothermal reaction at 1500 C. Y2O3 behaved as a catalyst to induce and favor the growth of glassy particles of YeAleSieOeC system at 1460e1540 C, of which the abnormal growth at 1560e1580 C became detrimental to the physical properties of the composites. (2) Al2O3 and SiC in situ formed from the carbothermal reduction of mullite could improve the heat capacity and the thermal conductivity of the composite ceramics, whilst the composite ceramics sintered at 1540 C exhibited the optimal physical and thermophysical properties, which were listed as follows: 46.7 MPa for bending strength, 11.0% for water absorption, 1.16 J/(g$K) for heat capacity, 13.73 W/(m$K) for thermal conductivity. The relatively high thermal conductivity endowed the composite ceramics with the good thermal shock resistance. (3) The advantages of the Al2O3/SiC composite ceramics for thermal storage applications were the lower sintering temperature, the simpler synthesis conditions, the lower manufacturing cost than SiC ceramic, and the higher thermal conductivity, the higher bending strength and the better thermal shock resistance compared with oxide ceramics. Acknowledgment Authors are very grateful to the financial support from “973 Program (2010CB227105)”, P. R. China. References [1] Dynamic simulation of thermal energy storage system of Badaling 1MW solar power tower plant, Renew. Energ 39 (2012) 455e462. [2] Y. Tian, C.Y. Zhao, A review of solar collectors and thermal energy storage in solar thermal applications, Appl. Energ 104 (2013) 538e553. [3] H. Singh, R.P. Saini, J.S. Saini, A review on packed bed solar energy storage systems, Renew. Sustain, Ener. Rev. 14 (2010) 1059e1069. , E. Oro , D. Boer, L.F. Cabeza, Embodied energy in thermal energy [4] L. Miro storage (TES) systems for high temperature applications, Appl. Energ 137 (2015) 793e799. [5] A.I. Fernandez, M. Martínez, M. Segarra, I. Martorell, L.F. Cabeza, Selection of materials with potential in sensible thermal energy storage, Sol. Energ. Mat. Sol. C 94 (2010) 1723e1729. [6] S. Kuravi, J. Trahan, D. Yogi Goswami, M.M. Rahman, E.K. Stefanakos, Thermal energy storage technologies and systems for concentrating solar power plants, Prog. Energ. Combust. 39 (2013) 285e319. , Eduard Oro , Dieter Boer, Luisa F. Cabeza, Embodied energy in [7] Laia Miro thermal energy storage (TES) systems for high temperature applications, Appl. Energ 137 (2015) 793e799. [8] S. Khare, M. Dell’Amico, C. Knight, S. McGarry, Selection of materials for high temperature sensible energy storage, Sol. Energ. Mat. Sol. C 115 (2013) 114e122. [9] G. Mebrahitom Asmelash, O. Mamat, F. Ahmad, A.K. Prasadarao, Thermal shock and fatigue behavior of pressureless sintered Al2O3-SiO2-ZrO2 composites, J. Advan. Ceram. 4 (2015) 190e198. [10] Y.X. Wang, M. Dong, Research development of the thermal shock resistance of thermal storage honeycomb ceramic, Chin. Ceram. 47 (2011) 1e6. [11] Y.J. Zhong, Q.G. Guo, S.Z. Li, J.L. Shi, L. Liu, Heat transfer enhancement of paraffin wax using carbon foam for thermal energy storage, Sol. Energ. Mat. Sol. C 94 (2010) 1011e1014. [12] G. Zanganeh, A. Pedretti, A. Haselbacher, A. Steinfeld, Design of packed bed thermal energy storage systems for high-temperature industrial process heat, Appl. Eng. 137 (2015) 812e822. [13] X.H. Xu, F. Zhao, J.F. Wu, J. Li, C.G. Li, Al2O3/SiC composite ceramics used in solar thermal power, J. Wuhan. Univ. Technol. 31 (13) (2009) 8e11. [14] R.X. Tan, Z.L. Ke, Y.C. Liu, S.A. Ze, H. Wang, P.A. Liu, J. Song, Porous ceramic materials used for heat storage and accumulation, J. Ceram. 2 (27) (2006) 217e226. [15] S.Q. Ding, Y.P. Zeng, D.L. Jiang, Thermal shock behavior of mullite-bonded porous silicon carbide ceramics with yttria addition, J. Phys. D. Appl. Phys. 40 (2007) 2138e2142.
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