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Effect of Si powder-supported catalyst on the microstructure and properties of Si3N4-MgO-C refractories
Construction and Building Materials 240 (2020) 117964
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Effect of Si powder-supported catalyst on the microstructure and properties of Si3N4-MgO-C refractories Yang Chen, Chengji Deng, Xing Wang, Jun Ding ⇑, Chao Yu, Hongxi Zhu The State Key Laboratory of Refractories and Metallurgy, Wuhan University of Science and Technology, Wuhan 430081, China
h i g h l i g h t s The effect of nitriding temperature on catalyst-free refractories was investigated. Catalytic nitridation method was used to control the growth of Si3N4 in refractory. Refractory with 1 wt% catalyst exhibited good properties and oxidation resistance. The oxidation resistance of refractory was studied based on TG and industrial CT.
a r t i c l e
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Article history: Received 16 July 2019 Received in revised form 23 December 2019 Accepted 25 December 2019
Keywords: Si3N4-MgO-C refractory b-Si3N4 Catalyst Oxidation resistance
a b s t r a c t Si3N4-MgO-C refractories were prepared using fused magnesia, flake graphite, and Si powder-supported Fe(NO3)39H2O as the raw materials and phenolic resin as the binder under the flow of N2. The effect of different nitriding temperatures on the morphology and production of Si3N4 in the uncatalyzed refractories was investigated. The effect of the morphology of b-Si3N4 on the properties of the refractories with different catalyst contents was investigated. The oxidation resistance of the refractories was also investigated. The results showed that the increase in the nitriding temperature was beneficial for the formation of b-Si3N4 in the refractories. However, very high nitriding temperatures generated a large number of pores in the refractories and deteriorated their properties. The refractory with 1 wt% catalyst exhibited good mechanical properties and oxidation resistance. However, the addition of excessive catalyst inhibited the formation of b-Si3N4 and deteriorated the performance of the refractories. Ó 2019 Elsevier Ltd. All rights reserved.
1. Introduction Carbon-containing refractories have been widely used for hightemperature applications, especially in the steel industry since the 1970s. The improvement in the preparation processes, microstructure, and performance of carbon-containing refractories and the enhancement of the service life of industrial furnaces (such as converters, ladle) have significantly reduced the production cost of the steel industry. Various carbon-containing refractories such as MgO-C [1], Al2O3-C [2–5], Al2O3-MgO-C [6–8], and self-bonded SiC refractories [9] have been reported. Among these, MgO-C refractories are important carbon composite refractories exhibiting excellent slag resistance, thermal shock resistance, and thermal conductivity. Moreover, the preparation process of MgO-C refractories is simple. These refractories are widely used as the lining
⇑ Corresponding author. E-mail address: [email protected] (J. Ding). https://doi.org/10.1016/j.conbuildmat.2019.117964 0950-0618/Ó 2019 Elsevier Ltd. All rights reserved.
material for electric furnaces, converters, and refining furnaces [10]. MgO-C refractories show excellent thermal shock resistance because of the high thermal conductivity and low elastic modulus of flake graphite. The low wettability of flake graphite and slag improves the corrosion resistance of MgO-C refractories. In addition, a carbon network structure is formed inside the material to improve the thermal shock resistance of MgO-C refractories. However, the layered structure of flake graphite makes it anisotropic, which tends to form layer cracks in the refractory during the manufacturing process and causes thermal stress concentration because of the mismatch in the thermal expansion during use. To overcome this limitation, artificial graphite materials such as nanocarbon black, carbon nanotubes, carbon nanofibers, and graphene have been used as alternatives to natural flake graphite. These artificial graphite materials absorb and reduce the thermal stress inside the refractory and toughen it [11–16]. Zhu et al. [10,17–19] investigated the effect of the use of artificial graphite materials such as expanded graphite, graphite oxide nanosheets,
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carbon nanotubes, and carbon black on the properties of MgO-C refractories. They found that the addition of these artificial graphite materials significantly improved the thermo-mechanical properties of MgO-C refractories. However, artificial graphite materials are expensive and their dispersion in MgO-C refractories is not uniform. Transition metal catalysts have been used in binders for refractories. These binders are then pyrolyzed and graphitized during the preparation process, followed by the in-situ generation of a onedimensional carbon structure [20,21]. Wei et al. [22] reported that the addition of Fe nanosheet-modified phenolic resin significantly improves the mechanical properties and thermal shock resistance of MgO-C refractories. This is because of the in-situ formation of carbon nanotubes, which create bridging and crack deflection in the refractory matrix. Rastegar et al. [23] used ferric nitratemodified phenolic resin as the binder to prepare a MgO-C refractory with 7 wt% modified phenolic resin. The use of the modified phenolic resin facilitated structural rearrangement in the matrix, thereby forming carbon nanotubes and a large number of ceramic
Table 1 Composition of the SMC refractories. Raw materials
Addition (wt%)
Fused magnesia (3–1 mm) Fused magnesia (1 mm) Silicon powder (0.074 mm) Flake graphite (0.15 mm) Fe(NO3)39H2O (addition) Liquid phenolic resin (addition)
30 40 20 10 0, 1, 3, 5 3.5
whiskers (MgAl2O4, MgO, Al4C3, and AlN), which improved the performance of the MgO-C refractory at high temperatures. However, like all carbon-containing refractories, MgO-C refractories suffer from the easy oxidization of carbon. The graphite present in MgO-C refractories oxidizes to form a decarburized layer. Moreover, the difference between the thermal expansion coefficients of MgO and graphite is large. These factors cause the loosening of the structure and a sharp decrease in its strength. Once the material is eroded and mechanically washed, the aggregate in the material erodes gradually, resulting in the layer-by-layer peeling off of the MgO-C refractory. Therefore, it is often necessary to add antioxidants such as Al, Si, TiO2, Al2O3, B4C, SiC, MgB2, Ti3AlC2, ZrSiO4, and Al4O4C to carbon-containing refractories to prevent the oxidation of the carbonaceous materials present in them [24–34]. The in-situ formation of a ceramic phase is also desirable. However, it is difficult to control the morphology and quantity of the in-situ formed ceramic phases. In our previous study [35], we added Si powder to a MgO-C refractory to prepare a Si3N4-MgO-C (SMC) refractory by nitriding. The b-Si3N4 phase so generated had a rod-like structure and was uniformly distributed in the material. It formed a direct interfacial chemical bond with MgO and graphite. This effectively increased the strength and toughness of the resulting refractory. However, Si3N4 in SMC refractories exists as both a-Si3N4 (particles, hexagonal plate, whisker) and b-Si3N4 (hexagonal columnar, cone). The morphology of Si3N4 significantly affects the properties of SMC refractories. Columnar Si3N4 can improve the properties of SMC refractories more effectively than granular Si3N4. Various studies have been carried out to prepare Si3N4 powders [36–40]. However, very few reports are available on the control-
Fig. 1. XRD patterns with magnified portion (2h = 20–40°) (a), and relative product contents (b) of the catalyst-free SMC refractories prepared at different nitriding temperatures.
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lable growth of Si3N4 formed in MgO-C refractories. In this study, we used catalytic nitridation to control the growth of Si3N4 in SMC refractories. The effect of the catalyst content on the formation of Si3N4 in SMC refractories was investigated systematically.
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In addition, the physical properties and oxidation resistance of the SMC refractories with different catalyst contents (0–5 wt%) were compared. Finally, the pore structure and distribution of the oxidized samples were investigated using industrial CT.
Fig. 2. SEM images of Si3N4 in the uncatalyzed SMC refractories nitrided at different temperatures: (a and b) 1350 °C; (c and d) 1400 °C; (e and f) 1450 °C; (g and h) 1500 °C.
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2. Experimental 2.1. Refractory preparation Fused magnesia (MgO, 3–1 mm and 1 mm, purity of 97.0 wt%, Gongyi Jinfeng Water Purification Material Co., Ltd., China), silicon powder (Si, 0.074 mm, purity of 99.9 wt%, Qinghe County Xintie Metal Co., Ltd., China), and flake graphite (C, 0.15 mm, purity of 97.7 wt%, Gongyi Jinfeng Water Purification Material Co., Ltd., China) were used as the raw materials. Ferric nitrate (Fe(NO3)39H2O, purity of 99.9 wt%, Shanghai Aladdin Biochemical Technology Co., Ltd., China) was used as the catalyst and phenolic resin (liquid, fixed carbon of 45.0–48.0 wt%, solid content of 75.5–78.5 wt% (150 °C), Jining Huakai Resin Co., Ltd., China) was used as the binder to prepare the SMC refractories. The composition of the SMC refractories is listed in Table 1. First, Si powder was pretreated and ferric nitrate was dissolved in absolute ethanol (C2H6O, content of 99.7 wt%, density of 0.789–0.791 g/mL (20 °C), Sinopharm Chemical Reagent Co., Ltd., China) in different amounts. Si powder was then added to the solution and the resulting mixture was stirred uniformly, followed by drying in an oven at 110 °C. According to the composition listed in Table 1, the mixed raw materials were pressed into strip samples with the dimensions of 25 mm 25 mm 140 mm and cylindrical samples with the dimensions of H36 mm U36 mm using a digital pressure testing machine (YES-2000, Tianshui Hongshan Testing Machine Co., Ltd., China) at 200 MPa. The samples were then cured at 110 °C for 24 h. The strip samples were used to evaluate the physical properties of the SMC refractories. The cylindrical samples on the other hand, were used to evaluate the oxidation resistance of the SMC refractories. The samples were placed in a high-temperature carbon tube sintering furnace (ZT-50–20, Shanghai Chen Hua Technology Co., Ltd., China) and were then heated from 25 to 1350 °C at a rate of 5 °C/min under the flow of nitrogen (nitrogen purity of 99.999%). The samples were held at this temperature for 2 h. Then, the samples were heated to 1400, 1450, and 1500 °C at 2 °C/min and held for 4 h. The sample fired at 1350 °C was directly heated to this temperature at a heating rate of 5 °C/min and was held at this temperature for 4 h. Finally, the fired samples were cooled to room temperature in the furnace.
nano-CT detection system (phoenix v|tome|x s-nanoCT) was used to examine the internal composition and structural distribution of the oxidized samples. 3. Results and discussion 3.1. Non-catalytic preparation of SMC refractories Fig. 1 shows the XRD patterns with magnified portion (2h = 20 40°), and relative product contents of the uncatalyzed samples nitrided at different temperatures. As can be observed from Fig. 1a, these samples consisted mainly of the a-Si3N4, b-Si3N4, SiC, MgO, graphite phases, and a small fraction of the Mg2SiO4 phase. The Si phase was present in the sample nitrided at 1350 °C, indicating that the Si in the raw material was not completely consumed during the reaction. With an increase in the nitriding temperature to 1400 °C, the intensity of the a-Si3N4 diffraction peak decreased slightly, the diffraction peak corresponding to elemental Si disappeared, and MgSiN2 could be detected. Since the melting point of elemental Si is 1410 °C, in the samples nitrided at temperatures higher than 1410 °C, Si existed in the liquid phase, while the a-Si3N4 phase rapidly transTable 2 b-Si3N4 crystal size of the uncatalyzed SMC refractories nitrided at different temperatures. Nitriding temperatures/°C
1400
1450
1500
Length/lm Thickness/lm
1.63–2.98 0.25–0.88
1.68–3.12 0.43–1.48
4.23–7.60 0.27–1.53
2.2. Characterization and testing The crystalline phases of the products were analyzed by X-ray diffraction (XRD, PANalytical, X’Pert Pro) using Cu Ka radiation (k = 1.540598 Å). Semi-quantitative analysis was carried out to calculate the relative contents of the samples. The morphology and crystal structure of the samples were observed by scanning electron microscopy (SEM, FEI, Nova 400 Nano SEM). The apparent porosity (AP) and bulk density (BD) of the samples were measured according to ISO 5017:2013 standard using a pore density analyzer (XQK-04, Wuxi Jianyi Experiment Instrument Co., Ltd., China). The cold modulus of rupture (CMOR) and cold crushing strength (CCS) of the samples were measured according to ISO 5014:1997 and ISO 8895:2004 standards using a microcomputer-controlled electronic universal testing machine (ETM1050, Shenzhen Wance Testing Machine Co., Ltd., China). Furthermore, the decarburized region of the nitrided samples subjected to firing in air was detected using Image-Pro Plus software. Prior to this, the oxidized samples were cut transversely along the vertical axis, and the cross-sectional images of the samples were captured using a digital camera [2,3]. The oxidation resistance of the samples was further analyzed using a thermal analyzer (STA449F3, NETZSCH). Finally, a high-precision micro-
Fig. 3. (a) BD and AP, (b) CMOR and CCS of the uncatalyzed SMC refractories nitrided at 1350–1500 °C.
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formed into the b-Si3N4 phase. This transformation can be attributed to structural reconstruction. Such a phase change usually requires the presence of a liquid phase and is completed by a dissolution-precipitation process [35,40]. As the nitriding temperature increased to 1450 °C, the intensity of the a-Si3N4 peak increased significantly. With a further increase in the nitriding temperature to 1500 °C, the intensity of the a-Si3N4 and b-Si3N4 diffraction peaks decreased, some of the a-Si3N4 diffraction peaks disappeared, and the intensity of the Mg2SiO4 diffraction peaks increased. The relative contents of the phases generated in the samples were calculated using Highscore Plus [41–43]. As can be observed from Fig. 1b, an increase in the nitriding temperature to 1350– 1400 °C facilitated the formation of Si3N4 and promoted the conversion of a-Si3N4 to b-Si3N4. The a-Si3N4, b-Si3N4, and MgSiN2 contents were the highest at 1400 °C. However, with a further increase in the nitriding temperature to 1500 °C, the a-Si3N4 and b-Si3N4 contents of the samples decreased along with a decrease in the MgSiN2 content to 0 wt% and an increase in the SiC and Mg2SiO4 contents. This can be attributed to the consumption of Si for the formation of Mg2SiO4 and SiC, thus suppressing the formation of Si3N4. Fig. 2 shows the morphologies of Si3N4 in the uncatalyzed samples nitrided at different temperatures. Because of the variation in the c-axis lattice constant, a-Si3N4 mainly shows particle-like and whisker-like morphology depending on the synthesis method [44]. However, the N atoms in b-Si3N4 grains exhibit different stacking speeds in the a-axis [2 1 0] and c-axis [0 0 1] directions. Thus, b-Si3N4 grains show lower interfacial energy and easier nucleation activation in the c-axis direction. The growth rate is higher in the
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c-axis direction. Hence, needle-like and cylindrical grains are formed in the c-axis direction [45]. As can be observed from Fig. 2a–h, hexagonal particle-like a-Si3N4 grains and hexagonal columnar b-Si3N4 grains were formed at high nitriding temperatures. At 1350 °C, closely packed a-Si3N4 grains were formed. These grains induced the dispersion toughening effect on the sample (Fig. 2a and b). At 1400–1450 °C, small columnar b-Si3N4 grains with different sizes were formed, and the b-Si3N4 growth followed the vapor–liquid-solid mechanism. The resulting b-Si3N4 grains were pinned inside the sample, which significantly increased the strength of the sample (Fig. 2c–f). At higher nitriding temperatures, a large number of long and thin hexagonal columnar b-Si3N4 grains were formed (Fig. 2g and h). In order to further analyze the effect of the nitriding temperature on the formation of Si3N4, Image-Pro Plus software was used to measure the grain size of b-Si3N4 formed at different temperatures (Table 2). With an increase in the nitriding temperature, the grain size of b-Si3N4 increased. Columnar b-Si3N4 with the length and thickness of 4.23–7.60 and 0.27–1.53 lm, respectively, showed the largest grain size at 1500 °C. This is because with an increase in the nitriding temperature, the grain growth accelerated and the viscosity of liquid Si decreased. The rate of dissolution, diffusion, and precipitation of a-Si3N4 in the liquid phase accelerated, which promoted the conversion of a-Si3N4 to b-Si3N4 and increased the grain size of b-Si3N4. Fig. 3 shows the physical properties of the uncatalyzed samples nitrided at different temperatures. As can be observed from Fig. 3a, the BD of the samples first increased with an increase in the nitriding temperature from 1350 to 1450 °C and then decreased with a further increase in the nitriding temperature to 1500 °C. The AP
Fig. 4. XRD patterns with magnified portion (2h = 20–40°) (a), and relative product contents (b) of the catalyzed SMC refractories (0–5 wt%) nitrided at 1400 °C.
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of the samples showed an opposing trend. At 1450 °C, the sample showed a large BD and a low AP of 2.76 g/cm3 and 21.54%, respectively. At 1500 °C, large amounts of various gases such as SiO, CO, and Mg were formed inside the sample. These gases generated a large number of pores and reduced the BD and increased the AP of the sample. Fig. 3b shows the CMOR and CCS curves of the samples without the catalyst nitrided at different temperatures. The sample nitrided at 1400 °C showed large CMOR and CCS values of 15.17 and 88.16 MPa, respectively. At 1450– 1500 °C, the CCS of the sample decreased significantly. As can be observed from Fig. 1b and 2, although the sample nitrided at
1500 °C showed large b-Si3N4 grains, it consisted of a small amount of b-Si3N4 and a large number of pores. This decreased the strength of the SMC refractory. The samples nitrided at 1400–1450 °C showed similar strengths, BD and AP values. The sample nitrided at 1450 °C (Fig. 3a) showed accelerated sintering and higher density than the sample nitrided at 1400 °C. However, at 1400 °C, a large number of ceramic phases (a-Si3N4, b-Si3N4, MgSiN2) were formed (Fig. 1b). The sample nitrided at 1400 °C showed pronounced reinforcement and toughening, and hence better mechanical properties than all the other samples (Fig. 3b). 3.2. Fe-catalyzed preparation of SMC refractories Fig. 4 shows the XRD patterns with magnified portion (2h = 20 40°), and relative product contents of the samples obtained using 0–5 wt% of the catalyst after nitriding at 1400 °C. As can be observed from Fig. 4a, the samples containing 0–5 wt% of the catalyst consisted of the a-Si3N4, b-Si3N4, SiC, MgO, graphite, and Mg2SiO4 phases. The sample containing 1 wt% of the catalyst showed lower a-Si3N4 diffraction peak intensity and higher b-Si3N4 diffraction peak intensity than the sample without the catalyst. Moreover, the sample with 1 wt% of the catalyst also showed unreacted Si. In the case of the sample with 3 wt% of the catalyst, the MgSiN2 phase disappeared. However, with a further increase in the catalyst content to 5 wt%, the intensity of the Si diffraction peak increased, indicating that the addition of too much catalyst was not conducive to the participation of Si in the reaction. In addition, the b-Si3N4 diffraction peaks disappeared partially. The semi-quantitative analysis (Fig. 4b) results showed that the catalyst-free sample prepared at 1400 °C produced 20 wt% a-Si3N4, 4 wt% b-Si3N4, 4 wt% SiC, 12 wt% Mg2SiO4, and 5 wt% MgSiN2.
Fig. 5. SEM images of Si3N4 in the catalyzed SMC refractories (1–5 wt%) nitrided at 1400 °C: (a) 1 wt%; (b) 3 wt%; (c) 5 wt%.
Fig. 6. (a) BD and AP, (b) CMOR and CCS of the catalyzed SMC refractories (0–5 wt%) nitrided at 1400 °C.
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When 1 wt% of the catalyst was added, the a-Si3N4 content of the sample reduced significantly to 10 wt%; the b-Si3N4 content increased to 6 wt%; 6 wt% SiC, 13 wt% Mg2SiO4, and 11 wt% MgSiN2 were formed. According to reaction (1), the Si3N4 participating in the reaction was a-Si3N4 and the addition of the catalyst increased the formation of b-Si3N4. With an increase in the catalyst content to 3 wt%, the a-Si3N4 and b-Si3N4 contents increased and the peak corresponding to MgSiN2 disappeared. With a further increase in the catalyst content to 5 wt%, the a-Si3N4 content increased and the b-Si3N4 content decreased, indicating that the addition of too much catalyst inhibited the conversion of the a-Si3N4 phase to the b-Si3N4 phase. This is because the addition of excessive catalyst causes the agglomeration and growth of particles at high temperatures, which ultimately results in a decrease in the catalytic activity or even the loss of catalytic properties [9].
3MgðgÞ þ Si3 N4ðsÞ þ N2ðgÞ ¼ 3MgSiN2ðsÞ
ð1Þ
Fig. 5 shows the SEM images of Si3N4 in the SMC refractories prepared using 1–5 wt% of the catalyst after nitriding at 1400 °C. As can be observed from Fig. 5a, the sample containing 1 wt% of the catalyst consisted of a large number of short columnar b-Si3N4 grains and a few hexagonal plate-like a-Si3N4 particles. The sample with 3 wt% of the catalyst showed irregular long columnar b-Si3N4 grains and small a-Si3N4 particles stacked into blocks (Fig. 5b). With an increase in the catalyst content to 5 wt %, a small number of short columnar b-Si3N4 grains were observed on flake graphite (Fig. 5c). The SEM images of the samples with different catalyst contents revealed that the catalyst content of 1 wt% was favorable for forming a good Si3N4 morphology inside the sample. Fig. 6 shows the physical properties of the samples nitrided at 1400 °C containing 0–5 wt% of the catalyst. With an increase in
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the catalyst content, the BD of the sample decreased, while the AP increased. However, the variation in the BD and AP values was not large. The BD value ranged from 2.62 to 2.68 g/cm3, while the AP value ranged from 21.99 to 25.00% (Fig. 6a). As can be observed from Fig. 6b, the addition of 1 wt% of the catalyst significantly improved the CMOR and CCS of the sample. However, an increase in the catalyst content to 3.0 wt% reduced the strength of the sample (especially the catalyst content of 3 wt%). Fig. 7a shows the anti-oxidation photographs of the samples (after nitriding at 1400 °C) fired at 1400 °C for 2 h in air containing 0–5 wt% of the catalyst. The sample containing 1 wt% of the catalyst showed the highest oxidation resistance, while the sample containing 5 wt% of the catalyst showed the lowest oxidation resistance. The decarburization area of Fig. 7a was calculated using Image-Pro Plus software and the results are shown in Fig. 7b. When 1 wt% of the catalyst was used, the sample showed a minimum oxidation index and its oxidation resistance improved. However, with an increase in the catalyst content, the oxidation index of the sample increased and the oxidation resistance decreased. Fig. 8 shows the XRD patterns with magnified portion (2h = 17 40°), and relative product contents of the oxidized decarburization areas of the samples shown in Fig. 7a. As can be observed from Fig. 8a, the samples with different catalyst contents showed the Mg2SiO4, SiO2, MgO, and MgSiO3 phases after oxidation. In the case of the samples with 0–3 wt% of the catalyst, Si2N2O was also formed and Si3N4 was not completely oxidized. Fig. 8b shows the relative contents of Si2N2O and Si3N4 in the oxidized decarburization areas of the samples. The samples containing 1 and 3 wt% of the catalyst showed large amounts of Si2N2O and Si3N4 after the oxidation in air. From Fig. 8b and 6a, it can be observed that the sample with 3 wt% of the catalyst showed the highest Si2N2O and Si3N4 contents, low BD, and high AP. Hence, the decarburization area of this sample was higher than that of
Fig. 7. Cross-sections (a) and oxidation indices (b) of the catalyzed SMC refractories (0–5 wt%) after the oxidation test at 1400 °C for 2 h in air.
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Fig. 8. XRD patterns with magnified portion (2h = 17–40°) (a), and relative product contents (b) of the oxidized decarburization area of the samples after firing at 1400 °C in an air atmosphere.
of the samples first increased, then decreased and finally increased. According to reactions (2–7), the oxidation of graphite reduced the mass of the samples, while the other reactions increased the mass of the samples. Therefore, the mass loss observed at 600–1000 °C can be attributed to the oxidation of graphite. The samples with 0, 1, 3, and 5 wt% of the catalyst showed 2, 2, 5, and 8 wt% mass loss, respectively. This result is consistent with those shown in Fig. 7b.
Fig. 9. TG curves of the catalyzed SMC refractories (0–5 wt%) after heating to 1400 °C at a rate of 10 °C/min in air.
the sample with 1 wt% of the catalyst. Si2N2O and Si3N4 were not present in the sample with 5 wt% of the catalyst, and the sample exhibited poor oxidation resistance. Thermogravimetric (TG) analysis was carried out in order to further investigate the effect of the catalyst content on the oxidation resistance of the SMC refractories (Fig. 9). The mass change
CðsÞ þ O2ðgÞ ¼ CO2ðgÞ
ð2Þ
SiðsÞ þ O2ðgÞ ¼ SiO2ðsÞ
ð3Þ
Si3 N4ðsÞ þ 3O2ðgÞ ¼ 3SiO2ðsÞ þ 2N2ðgÞ
ð4Þ
2Si3 N4ðsÞ þ 3O2ðgÞ ¼ 2Si2 N2 OðsÞ þ 2SiO2ðsÞ þ 2N2ðgÞ
ð5Þ
SiCðsÞ þ 2O2ðgÞ ¼ SiO2ðsÞ þ CO2ðgÞ
ð6Þ
2MgSiN2ðsÞ þ 3O2ðgÞ ¼ Mg2 SiO4ðsÞ þ SiO2ðsÞ þ 2N2ðgÞ
ð7Þ
High-resolution industrial CT is commonly used to scan samples to quantitatively characterize the three-dimensional spatial positions of their constituents [46,47]. The samples with the catalyst contents of 0 and 1 wt% (Fig. 7a) were tested and the results are shown in Figs. 10 and 11 (the gray part corresponds to the aggregate, the black part corresponds to the matrix, and the other colored parts surrounded by the white lines correspond to the closed pores). The catalyst-free sample underwent oxidation and
Y. Chen et al. / Construction and Building Materials 240 (2020) 117964
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Fig. 10. Industrial CT images of the uncatalyzed SMC refractory after firing at 1400 °C in air: (a) top view; (b) front view; (c) right view; (d) three-dimensional view.
showed pores with a diameter of 0.06–11.46 mm. The pores were mainly present between the aggregate and the matrix, and a part of the closed pores combined to form atmospheric pores (Fig. 10). The sample containing 1 wt% of the catalyst also underwent oxidation and showed pores with a diameter of 0.05–10.00 mm (Fig. 11). The material and closed pore volumes of the samples were calculated and the results are listed in Table 3. The closed porosity of the catalyst-free and 1 wt% catalyzed samples were 6.30% and 2.74%, respectively. This indicates that the sample containing 1 wt% of the catalyst showed better oxidation resistance than the catalystfree sample. 4. Conclusions Si3N4-MgO-C refractories were prepared by the in-situ nitridation of a Si powder-supported catalyst. The effect of the catalyst content on the microstructure, physical properties, and oxidation resistance of the Si3N4-MgO-C refractories were investigated. The catalyst-free sample exhibited better mechanical properties after nitriding at 1400 °C. Although the catalyst-free sample nitrided at 1500 °C showed crystallization and long columnar b-Si3N4 grains with a length of 4.23–7.60 lm and a thickness of 0.27–1.53 lm, the interior of the sample was porous, which severely deteriorated the properties of the sample. The addition of the catalyst facilitated the
conversion of a-Si3N4 to b-Si3N4 and promoted the formation of bSi3N4 in the MgO-C refractories. The refractory with 1 wt% of the catalyst showed significantly improved mechanical properties, oxidation resistance, and porosity as compared to the catalyst-free refractory.
CRediT authorship contribution statement Yang Chen: Conceptualization, Data curation, Writing - original draft, Writing - review & editing. Chengji Deng: Data curation, Writing - original draft, Writing - review & editing. Xing Wang: Data curation, Writing - original draft, Writing - review & editing. Jun Ding: Conceptualization, Data curation, Writing - original draft, Supervision, Writing - review & editing. Chao Yu: Data curation, Writing - original draft. Hongxi Zhu: Data curation, Writing original draft.
Declaration of Competing Interest The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.
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Fig. 11. Industrial CT images of the SMC refractory containing 1 wt% of the catalyst after firing at 1400 °C in air: (a) top view; (b) front view; (c) right view; (d) threedimensional view.
Table 3 Closed pore volumes of the samples after oxidation at 1400 °C. Catalyst content (wt%)
0
1
Material volume (mm3) Closed pore volume (mm3) Closed porosity (%)
704.47 47.38 6.30
909.73 25.63 2.74
Acknowledgments The authors acknowledge the financial support from the National Natural Science Foundation of China (grant nos. 51502215, 51574187 and 51602232) and the Natural Science Foundation of Hubei Province (grant no. 2018CFA022). References [1] X.C. Li, B.Q. Zhu, T.X. Wang, Effect of electromagnetic field on slag corrosion resistance of low carbon MgO-C refractories, Ceram. Int. 38 (2012) 2105–2109. [2] X. Wang, Y. Chen, C. Yu, J. Ding, D. Guo, C.J. Deng, H.X. Zhu, Preparation and application of ZrC-coated flake graphite for Al2O3-C refractories, J. Alloy. Compd. 788 (2019) 739–747. [3] Z.L. Liu, C.J. Deng, C. Yu, X. Wang, J. Ding, H.X. Zhu, Preparation of in situ grown silicon carbide whiskers onto graphite for application in Al2O3-C refractories, Ceram. Int. 44 (2018) 13944–13950.
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Construction and Building Materials journal homepage: www.elsevier.com/locate/conbuildmat
Effect of Si powder-supported catalyst on the microstructure and properties of Si3N4-MgO-C refractories Yang Chen, Chengji Deng, Xing Wang, Jun Ding ⇑, Chao Yu, Hongxi Zhu The State Key Laboratory of Refractories and Metallurgy, Wuhan University of Science and Technology, Wuhan 430081, China
h i g h l i g h t s The effect of nitriding temperature on catalyst-free refractories was investigated. Catalytic nitridation method was used to control the growth of Si3N4 in refractory. Refractory with 1 wt% catalyst exhibited good properties and oxidation resistance. The oxidation resistance of refractory was studied based on TG and industrial CT.
a r t i c l e
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Article history: Received 16 July 2019 Received in revised form 23 December 2019 Accepted 25 December 2019
Keywords: Si3N4-MgO-C refractory b-Si3N4 Catalyst Oxidation resistance
a b s t r a c t Si3N4-MgO-C refractories were prepared using fused magnesia, flake graphite, and Si powder-supported Fe(NO3)39H2O as the raw materials and phenolic resin as the binder under the flow of N2. The effect of different nitriding temperatures on the morphology and production of Si3N4 in the uncatalyzed refractories was investigated. The effect of the morphology of b-Si3N4 on the properties of the refractories with different catalyst contents was investigated. The oxidation resistance of the refractories was also investigated. The results showed that the increase in the nitriding temperature was beneficial for the formation of b-Si3N4 in the refractories. However, very high nitriding temperatures generated a large number of pores in the refractories and deteriorated their properties. The refractory with 1 wt% catalyst exhibited good mechanical properties and oxidation resistance. However, the addition of excessive catalyst inhibited the formation of b-Si3N4 and deteriorated the performance of the refractories. Ó 2019 Elsevier Ltd. All rights reserved.
1. Introduction Carbon-containing refractories have been widely used for hightemperature applications, especially in the steel industry since the 1970s. The improvement in the preparation processes, microstructure, and performance of carbon-containing refractories and the enhancement of the service life of industrial furnaces (such as converters, ladle) have significantly reduced the production cost of the steel industry. Various carbon-containing refractories such as MgO-C [1], Al2O3-C [2–5], Al2O3-MgO-C [6–8], and self-bonded SiC refractories [9] have been reported. Among these, MgO-C refractories are important carbon composite refractories exhibiting excellent slag resistance, thermal shock resistance, and thermal conductivity. Moreover, the preparation process of MgO-C refractories is simple. These refractories are widely used as the lining
⇑ Corresponding author. E-mail address: [email protected] (J. Ding). https://doi.org/10.1016/j.conbuildmat.2019.117964 0950-0618/Ó 2019 Elsevier Ltd. All rights reserved.
material for electric furnaces, converters, and refining furnaces [10]. MgO-C refractories show excellent thermal shock resistance because of the high thermal conductivity and low elastic modulus of flake graphite. The low wettability of flake graphite and slag improves the corrosion resistance of MgO-C refractories. In addition, a carbon network structure is formed inside the material to improve the thermal shock resistance of MgO-C refractories. However, the layered structure of flake graphite makes it anisotropic, which tends to form layer cracks in the refractory during the manufacturing process and causes thermal stress concentration because of the mismatch in the thermal expansion during use. To overcome this limitation, artificial graphite materials such as nanocarbon black, carbon nanotubes, carbon nanofibers, and graphene have been used as alternatives to natural flake graphite. These artificial graphite materials absorb and reduce the thermal stress inside the refractory and toughen it [11–16]. Zhu et al. [10,17–19] investigated the effect of the use of artificial graphite materials such as expanded graphite, graphite oxide nanosheets,
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carbon nanotubes, and carbon black on the properties of MgO-C refractories. They found that the addition of these artificial graphite materials significantly improved the thermo-mechanical properties of MgO-C refractories. However, artificial graphite materials are expensive and their dispersion in MgO-C refractories is not uniform. Transition metal catalysts have been used in binders for refractories. These binders are then pyrolyzed and graphitized during the preparation process, followed by the in-situ generation of a onedimensional carbon structure [20,21]. Wei et al. [22] reported that the addition of Fe nanosheet-modified phenolic resin significantly improves the mechanical properties and thermal shock resistance of MgO-C refractories. This is because of the in-situ formation of carbon nanotubes, which create bridging and crack deflection in the refractory matrix. Rastegar et al. [23] used ferric nitratemodified phenolic resin as the binder to prepare a MgO-C refractory with 7 wt% modified phenolic resin. The use of the modified phenolic resin facilitated structural rearrangement in the matrix, thereby forming carbon nanotubes and a large number of ceramic
Table 1 Composition of the SMC refractories. Raw materials
Addition (wt%)
Fused magnesia (3–1 mm) Fused magnesia (1 mm) Silicon powder (0.074 mm) Flake graphite (0.15 mm) Fe(NO3)39H2O (addition) Liquid phenolic resin (addition)
30 40 20 10 0, 1, 3, 5 3.5
whiskers (MgAl2O4, MgO, Al4C3, and AlN), which improved the performance of the MgO-C refractory at high temperatures. However, like all carbon-containing refractories, MgO-C refractories suffer from the easy oxidization of carbon. The graphite present in MgO-C refractories oxidizes to form a decarburized layer. Moreover, the difference between the thermal expansion coefficients of MgO and graphite is large. These factors cause the loosening of the structure and a sharp decrease in its strength. Once the material is eroded and mechanically washed, the aggregate in the material erodes gradually, resulting in the layer-by-layer peeling off of the MgO-C refractory. Therefore, it is often necessary to add antioxidants such as Al, Si, TiO2, Al2O3, B4C, SiC, MgB2, Ti3AlC2, ZrSiO4, and Al4O4C to carbon-containing refractories to prevent the oxidation of the carbonaceous materials present in them [24–34]. The in-situ formation of a ceramic phase is also desirable. However, it is difficult to control the morphology and quantity of the in-situ formed ceramic phases. In our previous study [35], we added Si powder to a MgO-C refractory to prepare a Si3N4-MgO-C (SMC) refractory by nitriding. The b-Si3N4 phase so generated had a rod-like structure and was uniformly distributed in the material. It formed a direct interfacial chemical bond with MgO and graphite. This effectively increased the strength and toughness of the resulting refractory. However, Si3N4 in SMC refractories exists as both a-Si3N4 (particles, hexagonal plate, whisker) and b-Si3N4 (hexagonal columnar, cone). The morphology of Si3N4 significantly affects the properties of SMC refractories. Columnar Si3N4 can improve the properties of SMC refractories more effectively than granular Si3N4. Various studies have been carried out to prepare Si3N4 powders [36–40]. However, very few reports are available on the control-
Fig. 1. XRD patterns with magnified portion (2h = 20–40°) (a), and relative product contents (b) of the catalyst-free SMC refractories prepared at different nitriding temperatures.
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lable growth of Si3N4 formed in MgO-C refractories. In this study, we used catalytic nitridation to control the growth of Si3N4 in SMC refractories. The effect of the catalyst content on the formation of Si3N4 in SMC refractories was investigated systematically.
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In addition, the physical properties and oxidation resistance of the SMC refractories with different catalyst contents (0–5 wt%) were compared. Finally, the pore structure and distribution of the oxidized samples were investigated using industrial CT.
Fig. 2. SEM images of Si3N4 in the uncatalyzed SMC refractories nitrided at different temperatures: (a and b) 1350 °C; (c and d) 1400 °C; (e and f) 1450 °C; (g and h) 1500 °C.
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2. Experimental 2.1. Refractory preparation Fused magnesia (MgO, 3–1 mm and 1 mm, purity of 97.0 wt%, Gongyi Jinfeng Water Purification Material Co., Ltd., China), silicon powder (Si, 0.074 mm, purity of 99.9 wt%, Qinghe County Xintie Metal Co., Ltd., China), and flake graphite (C, 0.15 mm, purity of 97.7 wt%, Gongyi Jinfeng Water Purification Material Co., Ltd., China) were used as the raw materials. Ferric nitrate (Fe(NO3)39H2O, purity of 99.9 wt%, Shanghai Aladdin Biochemical Technology Co., Ltd., China) was used as the catalyst and phenolic resin (liquid, fixed carbon of 45.0–48.0 wt%, solid content of 75.5–78.5 wt% (150 °C), Jining Huakai Resin Co., Ltd., China) was used as the binder to prepare the SMC refractories. The composition of the SMC refractories is listed in Table 1. First, Si powder was pretreated and ferric nitrate was dissolved in absolute ethanol (C2H6O, content of 99.7 wt%, density of 0.789–0.791 g/mL (20 °C), Sinopharm Chemical Reagent Co., Ltd., China) in different amounts. Si powder was then added to the solution and the resulting mixture was stirred uniformly, followed by drying in an oven at 110 °C. According to the composition listed in Table 1, the mixed raw materials were pressed into strip samples with the dimensions of 25 mm 25 mm 140 mm and cylindrical samples with the dimensions of H36 mm U36 mm using a digital pressure testing machine (YES-2000, Tianshui Hongshan Testing Machine Co., Ltd., China) at 200 MPa. The samples were then cured at 110 °C for 24 h. The strip samples were used to evaluate the physical properties of the SMC refractories. The cylindrical samples on the other hand, were used to evaluate the oxidation resistance of the SMC refractories. The samples were placed in a high-temperature carbon tube sintering furnace (ZT-50–20, Shanghai Chen Hua Technology Co., Ltd., China) and were then heated from 25 to 1350 °C at a rate of 5 °C/min under the flow of nitrogen (nitrogen purity of 99.999%). The samples were held at this temperature for 2 h. Then, the samples were heated to 1400, 1450, and 1500 °C at 2 °C/min and held for 4 h. The sample fired at 1350 °C was directly heated to this temperature at a heating rate of 5 °C/min and was held at this temperature for 4 h. Finally, the fired samples were cooled to room temperature in the furnace.
nano-CT detection system (phoenix v|tome|x s-nanoCT) was used to examine the internal composition and structural distribution of the oxidized samples. 3. Results and discussion 3.1. Non-catalytic preparation of SMC refractories Fig. 1 shows the XRD patterns with magnified portion (2h = 20 40°), and relative product contents of the uncatalyzed samples nitrided at different temperatures. As can be observed from Fig. 1a, these samples consisted mainly of the a-Si3N4, b-Si3N4, SiC, MgO, graphite phases, and a small fraction of the Mg2SiO4 phase. The Si phase was present in the sample nitrided at 1350 °C, indicating that the Si in the raw material was not completely consumed during the reaction. With an increase in the nitriding temperature to 1400 °C, the intensity of the a-Si3N4 diffraction peak decreased slightly, the diffraction peak corresponding to elemental Si disappeared, and MgSiN2 could be detected. Since the melting point of elemental Si is 1410 °C, in the samples nitrided at temperatures higher than 1410 °C, Si existed in the liquid phase, while the a-Si3N4 phase rapidly transTable 2 b-Si3N4 crystal size of the uncatalyzed SMC refractories nitrided at different temperatures. Nitriding temperatures/°C
1400
1450
1500
Length/lm Thickness/lm
1.63–2.98 0.25–0.88
1.68–3.12 0.43–1.48
4.23–7.60 0.27–1.53
2.2. Characterization and testing The crystalline phases of the products were analyzed by X-ray diffraction (XRD, PANalytical, X’Pert Pro) using Cu Ka radiation (k = 1.540598 Å). Semi-quantitative analysis was carried out to calculate the relative contents of the samples. The morphology and crystal structure of the samples were observed by scanning electron microscopy (SEM, FEI, Nova 400 Nano SEM). The apparent porosity (AP) and bulk density (BD) of the samples were measured according to ISO 5017:2013 standard using a pore density analyzer (XQK-04, Wuxi Jianyi Experiment Instrument Co., Ltd., China). The cold modulus of rupture (CMOR) and cold crushing strength (CCS) of the samples were measured according to ISO 5014:1997 and ISO 8895:2004 standards using a microcomputer-controlled electronic universal testing machine (ETM1050, Shenzhen Wance Testing Machine Co., Ltd., China). Furthermore, the decarburized region of the nitrided samples subjected to firing in air was detected using Image-Pro Plus software. Prior to this, the oxidized samples were cut transversely along the vertical axis, and the cross-sectional images of the samples were captured using a digital camera [2,3]. The oxidation resistance of the samples was further analyzed using a thermal analyzer (STA449F3, NETZSCH). Finally, a high-precision micro-
Fig. 3. (a) BD and AP, (b) CMOR and CCS of the uncatalyzed SMC refractories nitrided at 1350–1500 °C.
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formed into the b-Si3N4 phase. This transformation can be attributed to structural reconstruction. Such a phase change usually requires the presence of a liquid phase and is completed by a dissolution-precipitation process [35,40]. As the nitriding temperature increased to 1450 °C, the intensity of the a-Si3N4 peak increased significantly. With a further increase in the nitriding temperature to 1500 °C, the intensity of the a-Si3N4 and b-Si3N4 diffraction peaks decreased, some of the a-Si3N4 diffraction peaks disappeared, and the intensity of the Mg2SiO4 diffraction peaks increased. The relative contents of the phases generated in the samples were calculated using Highscore Plus [41–43]. As can be observed from Fig. 1b, an increase in the nitriding temperature to 1350– 1400 °C facilitated the formation of Si3N4 and promoted the conversion of a-Si3N4 to b-Si3N4. The a-Si3N4, b-Si3N4, and MgSiN2 contents were the highest at 1400 °C. However, with a further increase in the nitriding temperature to 1500 °C, the a-Si3N4 and b-Si3N4 contents of the samples decreased along with a decrease in the MgSiN2 content to 0 wt% and an increase in the SiC and Mg2SiO4 contents. This can be attributed to the consumption of Si for the formation of Mg2SiO4 and SiC, thus suppressing the formation of Si3N4. Fig. 2 shows the morphologies of Si3N4 in the uncatalyzed samples nitrided at different temperatures. Because of the variation in the c-axis lattice constant, a-Si3N4 mainly shows particle-like and whisker-like morphology depending on the synthesis method [44]. However, the N atoms in b-Si3N4 grains exhibit different stacking speeds in the a-axis [2 1 0] and c-axis [0 0 1] directions. Thus, b-Si3N4 grains show lower interfacial energy and easier nucleation activation in the c-axis direction. The growth rate is higher in the
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c-axis direction. Hence, needle-like and cylindrical grains are formed in the c-axis direction [45]. As can be observed from Fig. 2a–h, hexagonal particle-like a-Si3N4 grains and hexagonal columnar b-Si3N4 grains were formed at high nitriding temperatures. At 1350 °C, closely packed a-Si3N4 grains were formed. These grains induced the dispersion toughening effect on the sample (Fig. 2a and b). At 1400–1450 °C, small columnar b-Si3N4 grains with different sizes were formed, and the b-Si3N4 growth followed the vapor–liquid-solid mechanism. The resulting b-Si3N4 grains were pinned inside the sample, which significantly increased the strength of the sample (Fig. 2c–f). At higher nitriding temperatures, a large number of long and thin hexagonal columnar b-Si3N4 grains were formed (Fig. 2g and h). In order to further analyze the effect of the nitriding temperature on the formation of Si3N4, Image-Pro Plus software was used to measure the grain size of b-Si3N4 formed at different temperatures (Table 2). With an increase in the nitriding temperature, the grain size of b-Si3N4 increased. Columnar b-Si3N4 with the length and thickness of 4.23–7.60 and 0.27–1.53 lm, respectively, showed the largest grain size at 1500 °C. This is because with an increase in the nitriding temperature, the grain growth accelerated and the viscosity of liquid Si decreased. The rate of dissolution, diffusion, and precipitation of a-Si3N4 in the liquid phase accelerated, which promoted the conversion of a-Si3N4 to b-Si3N4 and increased the grain size of b-Si3N4. Fig. 3 shows the physical properties of the uncatalyzed samples nitrided at different temperatures. As can be observed from Fig. 3a, the BD of the samples first increased with an increase in the nitriding temperature from 1350 to 1450 °C and then decreased with a further increase in the nitriding temperature to 1500 °C. The AP
Fig. 4. XRD patterns with magnified portion (2h = 20–40°) (a), and relative product contents (b) of the catalyzed SMC refractories (0–5 wt%) nitrided at 1400 °C.
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of the samples showed an opposing trend. At 1450 °C, the sample showed a large BD and a low AP of 2.76 g/cm3 and 21.54%, respectively. At 1500 °C, large amounts of various gases such as SiO, CO, and Mg were formed inside the sample. These gases generated a large number of pores and reduced the BD and increased the AP of the sample. Fig. 3b shows the CMOR and CCS curves of the samples without the catalyst nitrided at different temperatures. The sample nitrided at 1400 °C showed large CMOR and CCS values of 15.17 and 88.16 MPa, respectively. At 1450– 1500 °C, the CCS of the sample decreased significantly. As can be observed from Fig. 1b and 2, although the sample nitrided at
1500 °C showed large b-Si3N4 grains, it consisted of a small amount of b-Si3N4 and a large number of pores. This decreased the strength of the SMC refractory. The samples nitrided at 1400–1450 °C showed similar strengths, BD and AP values. The sample nitrided at 1450 °C (Fig. 3a) showed accelerated sintering and higher density than the sample nitrided at 1400 °C. However, at 1400 °C, a large number of ceramic phases (a-Si3N4, b-Si3N4, MgSiN2) were formed (Fig. 1b). The sample nitrided at 1400 °C showed pronounced reinforcement and toughening, and hence better mechanical properties than all the other samples (Fig. 3b). 3.2. Fe-catalyzed preparation of SMC refractories Fig. 4 shows the XRD patterns with magnified portion (2h = 20 40°), and relative product contents of the samples obtained using 0–5 wt% of the catalyst after nitriding at 1400 °C. As can be observed from Fig. 4a, the samples containing 0–5 wt% of the catalyst consisted of the a-Si3N4, b-Si3N4, SiC, MgO, graphite, and Mg2SiO4 phases. The sample containing 1 wt% of the catalyst showed lower a-Si3N4 diffraction peak intensity and higher b-Si3N4 diffraction peak intensity than the sample without the catalyst. Moreover, the sample with 1 wt% of the catalyst also showed unreacted Si. In the case of the sample with 3 wt% of the catalyst, the MgSiN2 phase disappeared. However, with a further increase in the catalyst content to 5 wt%, the intensity of the Si diffraction peak increased, indicating that the addition of too much catalyst was not conducive to the participation of Si in the reaction. In addition, the b-Si3N4 diffraction peaks disappeared partially. The semi-quantitative analysis (Fig. 4b) results showed that the catalyst-free sample prepared at 1400 °C produced 20 wt% a-Si3N4, 4 wt% b-Si3N4, 4 wt% SiC, 12 wt% Mg2SiO4, and 5 wt% MgSiN2.
Fig. 5. SEM images of Si3N4 in the catalyzed SMC refractories (1–5 wt%) nitrided at 1400 °C: (a) 1 wt%; (b) 3 wt%; (c) 5 wt%.
Fig. 6. (a) BD and AP, (b) CMOR and CCS of the catalyzed SMC refractories (0–5 wt%) nitrided at 1400 °C.
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When 1 wt% of the catalyst was added, the a-Si3N4 content of the sample reduced significantly to 10 wt%; the b-Si3N4 content increased to 6 wt%; 6 wt% SiC, 13 wt% Mg2SiO4, and 11 wt% MgSiN2 were formed. According to reaction (1), the Si3N4 participating in the reaction was a-Si3N4 and the addition of the catalyst increased the formation of b-Si3N4. With an increase in the catalyst content to 3 wt%, the a-Si3N4 and b-Si3N4 contents increased and the peak corresponding to MgSiN2 disappeared. With a further increase in the catalyst content to 5 wt%, the a-Si3N4 content increased and the b-Si3N4 content decreased, indicating that the addition of too much catalyst inhibited the conversion of the a-Si3N4 phase to the b-Si3N4 phase. This is because the addition of excessive catalyst causes the agglomeration and growth of particles at high temperatures, which ultimately results in a decrease in the catalytic activity or even the loss of catalytic properties [9].
3MgðgÞ þ Si3 N4ðsÞ þ N2ðgÞ ¼ 3MgSiN2ðsÞ
ð1Þ
Fig. 5 shows the SEM images of Si3N4 in the SMC refractories prepared using 1–5 wt% of the catalyst after nitriding at 1400 °C. As can be observed from Fig. 5a, the sample containing 1 wt% of the catalyst consisted of a large number of short columnar b-Si3N4 grains and a few hexagonal plate-like a-Si3N4 particles. The sample with 3 wt% of the catalyst showed irregular long columnar b-Si3N4 grains and small a-Si3N4 particles stacked into blocks (Fig. 5b). With an increase in the catalyst content to 5 wt %, a small number of short columnar b-Si3N4 grains were observed on flake graphite (Fig. 5c). The SEM images of the samples with different catalyst contents revealed that the catalyst content of 1 wt% was favorable for forming a good Si3N4 morphology inside the sample. Fig. 6 shows the physical properties of the samples nitrided at 1400 °C containing 0–5 wt% of the catalyst. With an increase in
7
the catalyst content, the BD of the sample decreased, while the AP increased. However, the variation in the BD and AP values was not large. The BD value ranged from 2.62 to 2.68 g/cm3, while the AP value ranged from 21.99 to 25.00% (Fig. 6a). As can be observed from Fig. 6b, the addition of 1 wt% of the catalyst significantly improved the CMOR and CCS of the sample. However, an increase in the catalyst content to 3.0 wt% reduced the strength of the sample (especially the catalyst content of 3 wt%). Fig. 7a shows the anti-oxidation photographs of the samples (after nitriding at 1400 °C) fired at 1400 °C for 2 h in air containing 0–5 wt% of the catalyst. The sample containing 1 wt% of the catalyst showed the highest oxidation resistance, while the sample containing 5 wt% of the catalyst showed the lowest oxidation resistance. The decarburization area of Fig. 7a was calculated using Image-Pro Plus software and the results are shown in Fig. 7b. When 1 wt% of the catalyst was used, the sample showed a minimum oxidation index and its oxidation resistance improved. However, with an increase in the catalyst content, the oxidation index of the sample increased and the oxidation resistance decreased. Fig. 8 shows the XRD patterns with magnified portion (2h = 17 40°), and relative product contents of the oxidized decarburization areas of the samples shown in Fig. 7a. As can be observed from Fig. 8a, the samples with different catalyst contents showed the Mg2SiO4, SiO2, MgO, and MgSiO3 phases after oxidation. In the case of the samples with 0–3 wt% of the catalyst, Si2N2O was also formed and Si3N4 was not completely oxidized. Fig. 8b shows the relative contents of Si2N2O and Si3N4 in the oxidized decarburization areas of the samples. The samples containing 1 and 3 wt% of the catalyst showed large amounts of Si2N2O and Si3N4 after the oxidation in air. From Fig. 8b and 6a, it can be observed that the sample with 3 wt% of the catalyst showed the highest Si2N2O and Si3N4 contents, low BD, and high AP. Hence, the decarburization area of this sample was higher than that of
Fig. 7. Cross-sections (a) and oxidation indices (b) of the catalyzed SMC refractories (0–5 wt%) after the oxidation test at 1400 °C for 2 h in air.
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Fig. 8. XRD patterns with magnified portion (2h = 17–40°) (a), and relative product contents (b) of the oxidized decarburization area of the samples after firing at 1400 °C in an air atmosphere.
of the samples first increased, then decreased and finally increased. According to reactions (2–7), the oxidation of graphite reduced the mass of the samples, while the other reactions increased the mass of the samples. Therefore, the mass loss observed at 600–1000 °C can be attributed to the oxidation of graphite. The samples with 0, 1, 3, and 5 wt% of the catalyst showed 2, 2, 5, and 8 wt% mass loss, respectively. This result is consistent with those shown in Fig. 7b.
Fig. 9. TG curves of the catalyzed SMC refractories (0–5 wt%) after heating to 1400 °C at a rate of 10 °C/min in air.
the sample with 1 wt% of the catalyst. Si2N2O and Si3N4 were not present in the sample with 5 wt% of the catalyst, and the sample exhibited poor oxidation resistance. Thermogravimetric (TG) analysis was carried out in order to further investigate the effect of the catalyst content on the oxidation resistance of the SMC refractories (Fig. 9). The mass change
CðsÞ þ O2ðgÞ ¼ CO2ðgÞ
ð2Þ
SiðsÞ þ O2ðgÞ ¼ SiO2ðsÞ
ð3Þ
Si3 N4ðsÞ þ 3O2ðgÞ ¼ 3SiO2ðsÞ þ 2N2ðgÞ
ð4Þ
2Si3 N4ðsÞ þ 3O2ðgÞ ¼ 2Si2 N2 OðsÞ þ 2SiO2ðsÞ þ 2N2ðgÞ
ð5Þ
SiCðsÞ þ 2O2ðgÞ ¼ SiO2ðsÞ þ CO2ðgÞ
ð6Þ
2MgSiN2ðsÞ þ 3O2ðgÞ ¼ Mg2 SiO4ðsÞ þ SiO2ðsÞ þ 2N2ðgÞ
ð7Þ
High-resolution industrial CT is commonly used to scan samples to quantitatively characterize the three-dimensional spatial positions of their constituents [46,47]. The samples with the catalyst contents of 0 and 1 wt% (Fig. 7a) were tested and the results are shown in Figs. 10 and 11 (the gray part corresponds to the aggregate, the black part corresponds to the matrix, and the other colored parts surrounded by the white lines correspond to the closed pores). The catalyst-free sample underwent oxidation and
Y. Chen et al. / Construction and Building Materials 240 (2020) 117964
9
Fig. 10. Industrial CT images of the uncatalyzed SMC refractory after firing at 1400 °C in air: (a) top view; (b) front view; (c) right view; (d) three-dimensional view.
showed pores with a diameter of 0.06–11.46 mm. The pores were mainly present between the aggregate and the matrix, and a part of the closed pores combined to form atmospheric pores (Fig. 10). The sample containing 1 wt% of the catalyst also underwent oxidation and showed pores with a diameter of 0.05–10.00 mm (Fig. 11). The material and closed pore volumes of the samples were calculated and the results are listed in Table 3. The closed porosity of the catalyst-free and 1 wt% catalyzed samples were 6.30% and 2.74%, respectively. This indicates that the sample containing 1 wt% of the catalyst showed better oxidation resistance than the catalystfree sample. 4. Conclusions Si3N4-MgO-C refractories were prepared by the in-situ nitridation of a Si powder-supported catalyst. The effect of the catalyst content on the microstructure, physical properties, and oxidation resistance of the Si3N4-MgO-C refractories were investigated. The catalyst-free sample exhibited better mechanical properties after nitriding at 1400 °C. Although the catalyst-free sample nitrided at 1500 °C showed crystallization and long columnar b-Si3N4 grains with a length of 4.23–7.60 lm and a thickness of 0.27–1.53 lm, the interior of the sample was porous, which severely deteriorated the properties of the sample. The addition of the catalyst facilitated the
conversion of a-Si3N4 to b-Si3N4 and promoted the formation of bSi3N4 in the MgO-C refractories. The refractory with 1 wt% of the catalyst showed significantly improved mechanical properties, oxidation resistance, and porosity as compared to the catalyst-free refractory.
CRediT authorship contribution statement Yang Chen: Conceptualization, Data curation, Writing - original draft, Writing - review & editing. Chengji Deng: Data curation, Writing - original draft, Writing - review & editing. Xing Wang: Data curation, Writing - original draft, Writing - review & editing. Jun Ding: Conceptualization, Data curation, Writing - original draft, Supervision, Writing - review & editing. Chao Yu: Data curation, Writing - original draft. Hongxi Zhu: Data curation, Writing original draft.
Declaration of Competing Interest The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.
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Fig. 11. Industrial CT images of the SMC refractory containing 1 wt% of the catalyst after firing at 1400 °C in air: (a) top view; (b) front view; (c) right view; (d) threedimensional view.
Table 3 Closed pore volumes of the samples after oxidation at 1400 °C. Catalyst content (wt%)
0
1
Material volume (mm3) Closed pore volume (mm3) Closed porosity (%)
704.47 47.38 6.30
909.73 25.63 2.74
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