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Effect of carbon content on the oxidation resistance and kinetics of MgO-C refractory with the addition of Al powder
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Effect of carbon content on the oxidation resistance and kinetics of MgO-C refractory with the addition of Al powder Zhaoyang Liu, Jingkun Yu, Shuaijun Yue, Danbin Jia, Endong Jin, Beiyue Ma, Lei Yuan∗ School of Metallurgy, Northeastern University, Shenyang, 110819, PR China
ARTICLE INFO
ABSTRACT
Keywords: MgO-C refractory Indirect oxidation Oxidation resistance Oxidation kinetics
The effect of carbon content of MgO-C refractory was studied with respect to the physical properties, oxidation resistance and kinetics. The oxidation mechanism was investigated by calculating the relevant parameters of oxidation and analyzing the evolutions of phases and microstructures. The oxidation resistance of MgO-C refractory was influenced by the indirect oxidation of carbon through the formation of the MgO constituent layer. The effect of indirect oxidation resulted in the MgO-C refractory with 3 wt% carbon content having a lower oxidation rate and weight loss rate when fired at 1600 °C than when fired at 1400 °C. The oxidation kinetics results based on shrinking core model showed that the effective diffusion coefficient and full oxidation time increased with the increasing of carbon content. This seemingly paradoxical result was mainly due to the effect of the molar density of carbon prevailing over that of the effective diffusion coefficient in the oxidation process.
1. Introduction
and to maintain the function of carbon in refractories as long as possible, some antioxidants, such as metals (Al, Si, Mg, Fe), carbides (B4C, SiC) and oxides (TiO2, ZrO2) etc., are often applied to MgO-C refractory [7–10]. There are two mechanisms: one is the antioxidant is preferentially oxidized to the carbon, the other is the antioxidant forms a compound with carbon or other substances, blocking the pores formed by the carbon oxidation, which hinders the further inward diffusion of oxygen [11]. In addition, some measures such as coatings on the carbon surface and changing the carbon source have been implemented for improving the oxidation resistance of MgO-C refractories [12,13]. On the other hand, oxidation kinetics is an effective means to investigate the oxidation behaviour of MgO-C refractories from another point of view. Li et al. [14] studied the rate of carbon burnout through measuring the amount of CO converted into CO2 as a function of reaction time. They suggested that the inward diffusion rate of oxygen from the exterior atmosphere is predominant for the oxidation kinetics. Sadrnezhaad et al. [15] compared the oxidation kinetics between MgOC and MgO-C-Al refractory at 600oC–1300 °C based on a modified shrinking core model and corroborated that a mixed controlling mechanism governs the oxidation rate. Jansson et al. [16] studied the corrosion of MgO-C refractory in air, Ar, CO and Ar/CO gas atmospheres. They indicated that the reaction rate is directly dependent on the oxygen potential in the ambient atmosphere. However, the carbon content was fixed in the above-mentioned studies and only a few researches have paid attention to the effect of carbon content on the
Protecting carbon from oxidation has been always a formidable problem for carbon bonded magnesia (MgO-C) refractories [1]. It has attracted increasing interest in recent years because of the growing need for conventional high-carbon refractories to be replaced by lowcarbon refractories [2]. Conventional MgO-C refractories usually contain 12-18 wt% of carbon, which is associated with the main disadvantages of high carbon pick-up in molten steel, high heat loss and a considerable amount of COx gas emission [3]. However, the reduction of carbon content in a MgO-C refractory decreases the thermal conductivity and increases the wettability with slag, causing an increase in thermal stress and a decrease in corrosion resistance [4]. Meanwhile, the oxidized layer (decarburized layer) normally has poor strength and can be easily eroded away by the molten steel and slag [5]. These changes due to the oxidation of carbon damage the comprehensive properties of MgO-C refractory and thereby resulting in reduction of service life. Also, any degradation to a MgO-C refractory is likely to form large inclusions in the steel products and reduce their mechanical properties and corrosion resistance [6]. Therefore, choosing a MgO-C refractory with appropriate carbon content and improving its oxidation resistance are becoming more crucial with the development of advanced steelmaking technology and the requirement of high-quality steel products. In order to improve the oxidation resistance of MgO-C refractories
∗
Corresponding author. E-mail address: [email protected] (L. Yuan).
https://doi.org/10.1016/j.ceramint.2019.10.010 Received 6 September 2019; Received in revised form 20 September 2019; Accepted 2 October 2019 0272-8842/ © 2019 Elsevier Ltd and Techna Group S.r.l. All rights reserved.
Please cite this article as: Zhaoyang Liu, et al., Ceramics International, https://doi.org/10.1016/j.ceramint.2019.10.010
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furnace. The temperature schedule is shown in Fig. 2a, i.e. respectively heated to 1000 °C and 1400 °C at the rate of 10 °C·min−1 and soaking times were both 2 h. Oxidation tests were conducted in a muffle furnace and the detail of the heating process is shown in Fig. 2b, i.e. heated to 800 °C at the rate of 10 °C·min−1 firstly and then heated to the desired temperature at the rate of 5 °C·min−1 and soaking times were both 1 h. The oxidation kinetics of the MgO-C refractories was also tested by the muffle furnace. When the temperature stabilized at 1000 °C, the MgO-C refractories covered with two corundum plates top and bottom were put into the furnace and fired for different durations. After firing, the MgO-C refractories were taken out and cooled to room temperature in air. The bulk density and apparent porosity of the MgO-C refractories before and after coking were measured by the Archimedes method. The cold crushing strength was tested by an electronic material testing machine (WOW-100, Chuangbai Equipment Co., Ltd, Jinan, China) with the load rate of 0.5 mm min−1. The oxidized MgO-C refractories were cut along the longitudinal section and the thickness of the decarburized layer was measured, and the oxidation rate was calculated by Eq. (1). The masses of MgO-C refractories were weighed before and after oxidation, and the mass loss rate was calculated by Eq. (2). The phase compositions of the oxidized shell and the un-oxidized core were examined by X-ray diffractometry (XRD-7000, Shimadzu, Japan). The morphologies of the MgO-C refractories were observed by scanning electron microscopy (SEM, S-4800, Hitachi, Japan). The SEM samples were packed in epoxy resin before observation.
Table 1 Chemical compositions of MgO-C refractories (wt %). Raw materials
C1
C2
C3
C4
1-3 mm MgO 0.077-1 mm ≤0.077 mm Flake graphite Aluminium powder Phenolic resin
39 10 48 3 +2 +4
38 9 45 8 +2 +4
37 8 43 12 +2 +4
36 7 41 16 +2 +4
oxidation kinetics of MgO-C refractory. In this study, the oxidation resistance and mechanism of MgO-C refractories are discussed by analysing the phases and microstructures of the oxidized MgO-C refractories. Also, a shrinking core model was applied to study the effect of carbon content on the oxidation kinetics of MgO-C refractories, and the effective diffusion coefficients and full oxidation times were calculated. 2. Experimental MgO-C refractories with four different carbon contents were prepared and their chemical compositions are shown in Table 1. They are labelled as C1, C2, C3 and C4 according to their carbon content (3 wt%, 8 wt%, 12 wt% and 16 wt%). The main raw materials used were different sizes of magnesia (> 98 wt%), flake graphite (> 98 wt%), aluminium powder (analytical grade) and phenolic resin (liquid, 48% fixed carbon). Aluminium powder was purchased from Sinopharm Chemical Reagent Co., Ltd (Shanghai, China) and the others were purchased from Zhong Min Chi Yuan Industry Co., Ltd (Dashiqiao, China). The preparation process was as follows (Fig. 1): firstly, the coarse particles of MgO were uniformly mixed with the phenolic resin, and then the mixture of graphite, MgO powder, and aluminium powder was added to the mixture of MgO particles. The raw materials were mixed uniformly using a ball mill, and then placed in a mould with a diameter of 50 mm and pressed with a pressure of 200 MPa, maintained for 10 min. After pressing, the cylindrical MgO-C refractory (Φ50mm × ~40 mm) was cured at 200 °C for 24 h. After curing, the MgO-C refractories were embedded in graphite powder in an alumina crucible with a lid and heated in a resistance
ro =
V0
V V0
× 100% =
R02 h 0 R2h × 100% R 02 h 0
(1)
where ro is the oxidation rate (%), V0, R0 and h0 are the volume (m3), diameter (m) and height (m) of the MgO-C refractories before oxidation; V, R and h are the volume (m3), diameter (m) and height (m) of the unoxidized area following oxidation.
rm =
m0 m × 100% m0
(2)
where rm is the mass loss rate (%); m0 and m are the mass (g) before and after oxidation, respectively.
Fig. 1. Fabrication process of MgO-C refractories. 2
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Fig. 2. Temperature schedule of the coking (a) and oxidation (b) test.
Fig. 3. Bulk density (a) and apparent porosity (b) of cured and coked MgO-C refractories.
3. Results and discussion
oxidation rates of the MgO-C refractories were obtained by Eq. (1). The weight loss rates of MgO-C refractories were obtained by measuring the masses of MgO-C refractories before and after oxidation (Eq. (2)). The weight loss rate and oxidation rate results are shown in Fig. 6. From the figure, the weight loss rates increased while the oxidation rates decreased with increasing carbon content. Meanwhile, the weight loss rates and oxidation rates of MgO-C refractories fired at 1600 °C were higher than those fired at 1400 °C when the carbon content was 8 wt%, 12 wt%, and 16 wt%. However, abnormal results appeared when the carbon content was 3 wt%, and the weight loss rates and oxidation rates were lower at 1600 °C (marked by the red ellipse in Fig. 6). Due to the carbon content not changing the phase composition of MgO-C refractory, only the fired MgO-C refractory with 12 wt% carbon content was analysed in this study. The XRD patterns of the outer shell and the inner core are shown in Fig. 7. As can be seen from the figure, the main phases of the inner core were MgO, carbon, and MgAl2O4. A trace AlN phase and CaMgSiO4 impurity phase were also detected. The main phases of the outer shell were the same as in the inner core except without the carbon phase, which means that the carbon in the outer shell was completely oxidized. Meanwhile, the phase intensity of the MgAl2O4 was obviously increased in the outer shell samples. Aluminium powder reacted with nitrogen in the air to form AlN at high temperature (Eq. (3)). Also, part of the aluminium powder reacted with oxygen to form Al2O3 (Eq. (4)). The Al2O3 phase could also have been formed by the reaction between Al4C3 and CO (Eqs. (5) and (6)). The resulting alumina and magnesia formed MgAl2O4 (Eq. (7)).
3.1. Physical properties Fig. 3 shows the bulk density and apparent porosity of the MgO-C refractories with different carbon contents before and after coking. The bulk density of the cured MgO-C refractories decreased with increasing carbon content, which was mainly due to the lighter density of carbon than that of MgO. Meanwhile, the increasing of carbon content decreased the pores between MgO particles. After coking, the bulk density was decreased and the apparent porosity was increased significantly with the decomposing of phonetic resin to CH4, H2O, CO, C2H6, and H2 etc [17,18]. Different to the cured samples, the apparent porosity of the coked samples increased with the increasing of carbon content. Another noteworthy point is that the higher density and lower porosity were obtained in the 1400 °C batch. Fig. 4 shows the force-displacement curves and cold crushing strengths of MgO-C refractories before and after coking. From Fig. 4a–c, the failure displacements were decreased with the increasing of carbon content at 200 °C and 1000 °C, while it became ruleless when the coking temperature reached 1400 °C. From Fig. 4d, the cold crushing strength of the cured and coked MgO-C refractories decreased as the carbon content increased. Compared with the cured MgO-C refractories, the cold crushing strength of the coked MgO-C refractories were greatly reduced, and the higher the carbon content, the greater the reduction. 3.2. Oxidation resistance and mechanism The longitudinal section views of the fired MgO-C refractories are shown in Fig. 5. It can be seen that the MgO-C refractories have a distinct oxidized outer shell and an unoxidized inner core. Volumes of the shells and the cores were calculated from the above figure, and the 3
2Al(l, s)+N2(g)→2AlN(s)
(3)
4Al(l, s)+3O2(g)→2Al2O3(s)
(4)
4Al(l, s)+3C(s)→Al4C3(s)
(5)
Al4C3(s)+6CO(g)→2Al2O3(s)+9C(s)
(6)
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Fig. 4. Force-displacement curves: (a) cured at 200 °C, (b) coked at 1000 °C, (c) coked at 1400 °C and (d) cold crushing strength of MgO-C refractories.
Al2O3(s)+MgO(s)→MgAl2O4(s)
(7)
When oxidized at 1400 °C, the carbon in the MgO-C refractory reacted with oxygen to form carbon monoxide (direct oxidation, Eq. (8)), which formed the porous structure and isolated the MgO particles from each other in the decarburized layer (Fig. 8a-b). However, when the temperature was exceeded 1400 °C, the reaction between carbon and MgO started, which is the so-called indirect oxidation of carbon (Eq. (9)) [19]. The oxidation behaviour of the MgO-C refractory was the combined result of the direct oxidation and indirect oxidation of carbon. The reaction between MgO and carbon generated magnesium vapor, which was outwardly diffused through the pores of the decarburized layer where it reacted with oxygen to form MgO (Eq. (10)). The regenerated MgO partially filled the pores of the decarburized layer and connected the original MgO particles to form the continuous layer (Fig. 8c-d), which hindered the entry of the oxygen and improved the oxidation resistance of the MgO-C refractory to a certain extent [20]. The effect of indirect oxidation was more pronounced with lower
Fig. 6. Effect of carbon content on the oxidation rate and weight loss rate of MgO-C refractory.
Fig. 5. Longitudinal section view of the fired MgO-C refractories: (a) 1400 °C, (b) 1600 °C. 4
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Fig. 9. Effect of carbon content on the cold crushing strengths of fired MgO-C refractories. Fig. 7. XRD patterns of the fired MgO-C refractories: (a) 1400 °C, shell, (b) 1600 °C, shell, (c) 1400 °C, core, (d) 1600 °C, core.
is clear from Fig. 9 that, the cold crushing strengths fired at 1600 °C were higher than those fired at 1400 °C, especially for the fired MgO-C refractory with 3 wt% carbon content, which dramatically increased from 22.8 MPa to 56.5 MPa when the temperature increased from 1400 °C to 1600 °C. Undoubtedly, denser decarburized layer provided a better protection for the MgO-C refractory from the eroding away by the molten steel and slag [23].
carbon content in the MgO-C refractory, due to the MgO continuous layer being more easily formed when fewer pores were generated by the oxidation of carbon. Therefore, the MgO-C refractory with 3 wt% carbon content fired at 1600 °C had a lower weight loss rate and oxidation rate than that of fired at 1400 °C, even though the higher temperature means a faster reaction rate and mass transfer rate in general [21]. When the carbon content reached 8 wt% or higher, a more porous structure was created after oxidation of carbon, and the regenerated MgO was not sufficient to form a MgO continuous layer. Therefore, the effect of indirect oxidation was not so obvious and the weight loss rate and oxidation rate of the MgO-C refractories fired at 1600 °C were higher than those fired at 1400 °C. These results are in a good agreement with results of others [22]. 2C(s)+O2(g)→2CO(g)
(8)
C(s)+MgO(s)→CO(g)+Mg(g)
(9)
2 Mg(g)+O2(g)→MgO(s)
3.3. Oxidation kinetics Fig. 10 shows the cross section views of the MgO-C refractories after firing at 1000 °C for different durations. It can clearly be seen from the figure that the boundary of the non-oxidized core and the oxidized exterior shell are evident in four different carbon contents of MgO-C refractories. Meanwhile, the thicknesses of the oxidized exterior shells increased with the prolonging of firing time and decreased with the increasing of carbon content. According to the shrinking core model [24], the oxidation of MgO-C refractory can be divided into five parts: (Ⅰ) external mass transport of oxygen from the air flow bulk to the outer face of the sample, (Ⅱ) diffusion of oxygen through the pores in the decarbonized layer to reach the reaction interface, (Ⅲ) chemical reaction between graphite and oxygen at the reaction interface to form carbon monoxide, (Ⅳ)
(10)
The cold crushing strengths of the fired MgO-C refractories were also influenced by the formation of the MgO continuous layer (Fig. 9). It
Fig. 8. SEM images of the oxidized shell of fired MgO-C refractories: (a) 1400 °C, 3 wt% C, (b) 1400 °C, 16 wt% C, (c) 1600 °C, 3 wt% C, (d) 1600 °C, 16 wt% C. 5
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Fig. 10. Cross section views of fired MgO-C refractories: (a) 3 wt% C, (b) 8 wt% C, (c) 12 wt% C, (d) 16 wt% C.
diffusion of carbon monoxide through the decarbonized layer to the outer face, (Ⅴ) mass transfer of carbon monoxide from the outer face into the surrounding atmosphere. The schematic diagram of the MgO-C refractory oxidation process is shown in Fig. 11. When the temperature exceeded 800 °C, the external mass transport process of gas (parts Ⅰ and Ⅳ) and the reaction rate at the interface (part Ⅴ) were fast. Therefore, the gas diffusion in the decarburized layer can be deemed a restrictive factor and the oxidation rate equation of carbon can be shown as Eqs. (11) and (12) [25].
xC = 1
t=
rt r0
2
2 C r0 [x C + (1 4bDeff Cb
(11)
x C)ln(1
x C)]
(12)
where, t is the oxidation time (min), xC is the fractional oxidation of carbon at time t (%), ρC is the molar density of carbon in the MgO-C refractory (mol·cm−3), r0 is the initial radius of the MgO-C refractory (cm), b refers to the stoichiometric coefficient of carbon in the oxidation reaction with 1 mol of O2, Deff refers to the effective diffusion coefficient through the decarburized layer (cm2·min−1), and Cb is the oxygen concentration near the surface of the sample (mol·cm−3). According to the thicknesses of the decarburized layers measured
Fig. 11. Schematic diagram of the MgO-C refractory oxidation process.
6
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Fig. 12. Relationship between xC+(1-xC)ln(1-xC) and oxidation time. Fig. 13. Effect of carbon content on the effective diffusion coefficient and full oxidation times of MgO-C refractories.
Table 2 Related parameters of the oxidation rate equation for MgO-C refractories.
ρC × 10−2 (mol ⋅cm−3) r0 (mm) Cb × 10−6 (mol ⋅cm−3) b (-)
C1
C2
C3
C4
0.69 25 1.985 1.39
1.82 25 1.985 1.73
2.72 25 1.985 1.81
3.51 25 1.985 1.86
Table 3 Effective diffusion coefficient for previous oxidation tests.
from Fig. 10, the relationship between xC+(1-xC)ln(1-xC) and oxidation time can be obtained and is shown in Fig. 12. It can be seen that the two terms exhibits a good linear relation in the four different carbon contents of MgO-C refractories. The straight lines were fitted according to the dots and the slops were obtained. According to the related parameters of the oxidation rate equation shown in Table 2 and the definition of slope (Eq. (13)), the effective diffusion coefficient can be obtained. Meanwhile, the value of xC tended to the number 1 when the MgO-C refractory tended to totally oxidization, and the value of full oxidation time (tf) equalled to the reciprocal of the slope value (Eq. (14)).
4bDeff Cb S= 2 Cr0
tf =
2 1 C r0 = S 4bDeff Cb
Carbon content (wt %)
Temperature (oC)
Deff (cm2·min−1)
Reference number (-)
15 5 10 15 20 7 14 20
1000 1300
15.4 ~12 ~21 ~27 ~32 31.8 22.2 35.12 46.19 66.47 10.69–13.25 13.68–17.33
[14] [24]
14.3
1000 1000 1100 1200 950 1100
[26] [27] [28]
coefficients increased while the decarburized layer thicknesses decreased with the increasing of carbon content. This seemingly paradoxical result can be explained as follows. For a given decarburized layer thickness, the ratio of oxidation time (t) between x wt% carbon content (x is 3, 8, 12, 16) and 3 wt% carbon content of MgO-C refractory can be expressed as Eq. (16) [25]. The t was mainly related to the molar density of carbon (ρC) and effective diffusion coefficients (Deff) in the MgO-C refractories. Variations of ρC(x wt%)/ρC(3 wt%) and
(13) (14)
where, S is the slope of the straight line fitted by the experimental dots, tf is the full oxidation time (min). The results of effective diffusion coefficient and full oxidation time versus carbon content are shown in Fig. 13. It can be seen that the effective diffusion coefficient increased with the increasing of carbon content, which was mainly because the porosity of the decarburized layer was increased with the oxidation of carbon. This effect was partially impeded due to the reducing effect of carbon content on initial porosity of the MgO-C refractories (as seen in Fig. 3b). Meanwhile, the values of effective diffusion coefficients obtained in this study were much lower than that of others’ results (Table 3), which can be explained by the following: (1) 2 wt% of aluminium powder was added as antioxidant. The effect of aluminium powder on oxidation resistance of MgO-C refractory was not only on the preferential oxidization of aluminium over carbon, but also expansion in the specific volume accompanied by the formation of MgAl2O4 phase [29]. The expanded volume resulted in partial filling of the micro-pores of the decarburized layer and decreased inter diffusion rate of the gases through the solid and porous regions. (2) The hearth of the muffle furnace was a relatively closed space, where the mobility of the air around the MgO-C refractory was limited. As can be seen from Figs. 13 and Fig. 10, the effective diffusion
Fig. 14. Molar density ratio, effective diffusion coefficient ratio and full oxidation time ratio versus carbon content. 7
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Deff(x wt%)/Deff(3 wt%) with carbon content are shown in Fig. 14. It can be seen that both ratios increased with the increasing of carbon content. However, the increasing rate of the molar density ratio was higher than that of the effective diffusion coefficient ratio. Therefore, the oxidation time ratio increased with the increasing of carbon content. In other word, to form the same thickness of decarburized layer, longer oxidation time was required for MgO-C refractories with higher carbon content. For example, when the MgO-C refractory full oxidized (decarburized layer thickness of 25 mm), the full oxidation time was increased from 1339 min to 4566 min when the carbon content increased from 3 wt% to 16 wt% (Fig. 13). The increasing rate of the full oxidation time ratio was fell in between the increasing rate of the molar density ratio and the effective diffusion coefficient ratio (Fig. 14).
t (x wt%) C (x wt%)/ C (3 wt%) = t (3 wt%) Deff (x wt%)/ Deff (3 wt%)
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4. Conclusions Cured and coked MgO-C refractories were studied with respect to bulk density, apparent porosity, and cold crushing strength. The oxidation resistance and mechanism of MgO-C refractories were investigated via phase and microstructure evaluation. Also, the shrinking core model was applied in the oxidation kinetics of MgO-C refractories with different carbon content. Based on the above results, the following conclusions are drawn: (1) Bulk density, apparent porosity, and cold crushing strength of the cured MgO-C refractories decreased with the increasing of carbon content. Compared with the cured samples, the bulk density and cold crushing strength greatly reduced, and the apparent porosity significantly increased after curing at 1000 °C and 1400 °C. (2) Mg vapor was generated by the reaction between MgO and carbon at 1600 °C, which re-oxidized to MgO in the outward diffusion process of the decarburized layer. The regenerated MgO resulted in apparent porosity reduction and pore-filling effects, which caused diminution of the internal diffusion of gases in the decarburized layer. Finally, the oxidation resistance of MgO-C refractory with 3 wt% carbon content fired at 1600 °C was higher than that fired 1400 °C. (3) Effective diffusion coefficient increased while the thickness of the decarburized layer decreased with the increasing of carbon content. The seemingly paradoxical result was mainly due to the effect of the molar density of carbon prevailing over the effective diffusion coefficient in the oxidation process. As a result, longer oxidation time was required for full oxidation of MgO-C refractories with higher carbon content. 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. Acknowledgements The authors would like to express their gratitude for the financial support from the National Natural Science Foundation of China (Grant No. 51974074, 51874083 and 51404056) and the Fundamental
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Effect of carbon content on the oxidation resistance and kinetics of MgO-C refractory with the addition of Al powder Zhaoyang Liu, Jingkun Yu, Shuaijun Yue, Danbin Jia, Endong Jin, Beiyue Ma, Lei Yuan∗ School of Metallurgy, Northeastern University, Shenyang, 110819, PR China
ARTICLE INFO
ABSTRACT
Keywords: MgO-C refractory Indirect oxidation Oxidation resistance Oxidation kinetics
The effect of carbon content of MgO-C refractory was studied with respect to the physical properties, oxidation resistance and kinetics. The oxidation mechanism was investigated by calculating the relevant parameters of oxidation and analyzing the evolutions of phases and microstructures. The oxidation resistance of MgO-C refractory was influenced by the indirect oxidation of carbon through the formation of the MgO constituent layer. The effect of indirect oxidation resulted in the MgO-C refractory with 3 wt% carbon content having a lower oxidation rate and weight loss rate when fired at 1600 °C than when fired at 1400 °C. The oxidation kinetics results based on shrinking core model showed that the effective diffusion coefficient and full oxidation time increased with the increasing of carbon content. This seemingly paradoxical result was mainly due to the effect of the molar density of carbon prevailing over that of the effective diffusion coefficient in the oxidation process.
1. Introduction
and to maintain the function of carbon in refractories as long as possible, some antioxidants, such as metals (Al, Si, Mg, Fe), carbides (B4C, SiC) and oxides (TiO2, ZrO2) etc., are often applied to MgO-C refractory [7–10]. There are two mechanisms: one is the antioxidant is preferentially oxidized to the carbon, the other is the antioxidant forms a compound with carbon or other substances, blocking the pores formed by the carbon oxidation, which hinders the further inward diffusion of oxygen [11]. In addition, some measures such as coatings on the carbon surface and changing the carbon source have been implemented for improving the oxidation resistance of MgO-C refractories [12,13]. On the other hand, oxidation kinetics is an effective means to investigate the oxidation behaviour of MgO-C refractories from another point of view. Li et al. [14] studied the rate of carbon burnout through measuring the amount of CO converted into CO2 as a function of reaction time. They suggested that the inward diffusion rate of oxygen from the exterior atmosphere is predominant for the oxidation kinetics. Sadrnezhaad et al. [15] compared the oxidation kinetics between MgOC and MgO-C-Al refractory at 600oC–1300 °C based on a modified shrinking core model and corroborated that a mixed controlling mechanism governs the oxidation rate. Jansson et al. [16] studied the corrosion of MgO-C refractory in air, Ar, CO and Ar/CO gas atmospheres. They indicated that the reaction rate is directly dependent on the oxygen potential in the ambient atmosphere. However, the carbon content was fixed in the above-mentioned studies and only a few researches have paid attention to the effect of carbon content on the
Protecting carbon from oxidation has been always a formidable problem for carbon bonded magnesia (MgO-C) refractories [1]. It has attracted increasing interest in recent years because of the growing need for conventional high-carbon refractories to be replaced by lowcarbon refractories [2]. Conventional MgO-C refractories usually contain 12-18 wt% of carbon, which is associated with the main disadvantages of high carbon pick-up in molten steel, high heat loss and a considerable amount of COx gas emission [3]. However, the reduction of carbon content in a MgO-C refractory decreases the thermal conductivity and increases the wettability with slag, causing an increase in thermal stress and a decrease in corrosion resistance [4]. Meanwhile, the oxidized layer (decarburized layer) normally has poor strength and can be easily eroded away by the molten steel and slag [5]. These changes due to the oxidation of carbon damage the comprehensive properties of MgO-C refractory and thereby resulting in reduction of service life. Also, any degradation to a MgO-C refractory is likely to form large inclusions in the steel products and reduce their mechanical properties and corrosion resistance [6]. Therefore, choosing a MgO-C refractory with appropriate carbon content and improving its oxidation resistance are becoming more crucial with the development of advanced steelmaking technology and the requirement of high-quality steel products. In order to improve the oxidation resistance of MgO-C refractories
∗
Corresponding author. E-mail address: [email protected] (L. Yuan).
https://doi.org/10.1016/j.ceramint.2019.10.010 Received 6 September 2019; Received in revised form 20 September 2019; Accepted 2 October 2019 0272-8842/ © 2019 Elsevier Ltd and Techna Group S.r.l. All rights reserved.
Please cite this article as: Zhaoyang Liu, et al., Ceramics International, https://doi.org/10.1016/j.ceramint.2019.10.010
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furnace. The temperature schedule is shown in Fig. 2a, i.e. respectively heated to 1000 °C and 1400 °C at the rate of 10 °C·min−1 and soaking times were both 2 h. Oxidation tests were conducted in a muffle furnace and the detail of the heating process is shown in Fig. 2b, i.e. heated to 800 °C at the rate of 10 °C·min−1 firstly and then heated to the desired temperature at the rate of 5 °C·min−1 and soaking times were both 1 h. The oxidation kinetics of the MgO-C refractories was also tested by the muffle furnace. When the temperature stabilized at 1000 °C, the MgO-C refractories covered with two corundum plates top and bottom were put into the furnace and fired for different durations. After firing, the MgO-C refractories were taken out and cooled to room temperature in air. The bulk density and apparent porosity of the MgO-C refractories before and after coking were measured by the Archimedes method. The cold crushing strength was tested by an electronic material testing machine (WOW-100, Chuangbai Equipment Co., Ltd, Jinan, China) with the load rate of 0.5 mm min−1. The oxidized MgO-C refractories were cut along the longitudinal section and the thickness of the decarburized layer was measured, and the oxidation rate was calculated by Eq. (1). The masses of MgO-C refractories were weighed before and after oxidation, and the mass loss rate was calculated by Eq. (2). The phase compositions of the oxidized shell and the un-oxidized core were examined by X-ray diffractometry (XRD-7000, Shimadzu, Japan). The morphologies of the MgO-C refractories were observed by scanning electron microscopy (SEM, S-4800, Hitachi, Japan). The SEM samples were packed in epoxy resin before observation.
Table 1 Chemical compositions of MgO-C refractories (wt %). Raw materials
C1
C2
C3
C4
1-3 mm MgO 0.077-1 mm ≤0.077 mm Flake graphite Aluminium powder Phenolic resin
39 10 48 3 +2 +4
38 9 45 8 +2 +4
37 8 43 12 +2 +4
36 7 41 16 +2 +4
oxidation kinetics of MgO-C refractory. In this study, the oxidation resistance and mechanism of MgO-C refractories are discussed by analysing the phases and microstructures of the oxidized MgO-C refractories. Also, a shrinking core model was applied to study the effect of carbon content on the oxidation kinetics of MgO-C refractories, and the effective diffusion coefficients and full oxidation times were calculated. 2. Experimental MgO-C refractories with four different carbon contents were prepared and their chemical compositions are shown in Table 1. They are labelled as C1, C2, C3 and C4 according to their carbon content (3 wt%, 8 wt%, 12 wt% and 16 wt%). The main raw materials used were different sizes of magnesia (> 98 wt%), flake graphite (> 98 wt%), aluminium powder (analytical grade) and phenolic resin (liquid, 48% fixed carbon). Aluminium powder was purchased from Sinopharm Chemical Reagent Co., Ltd (Shanghai, China) and the others were purchased from Zhong Min Chi Yuan Industry Co., Ltd (Dashiqiao, China). The preparation process was as follows (Fig. 1): firstly, the coarse particles of MgO were uniformly mixed with the phenolic resin, and then the mixture of graphite, MgO powder, and aluminium powder was added to the mixture of MgO particles. The raw materials were mixed uniformly using a ball mill, and then placed in a mould with a diameter of 50 mm and pressed with a pressure of 200 MPa, maintained for 10 min. After pressing, the cylindrical MgO-C refractory (Φ50mm × ~40 mm) was cured at 200 °C for 24 h. After curing, the MgO-C refractories were embedded in graphite powder in an alumina crucible with a lid and heated in a resistance
ro =
V0
V V0
× 100% =
R02 h 0 R2h × 100% R 02 h 0
(1)
where ro is the oxidation rate (%), V0, R0 and h0 are the volume (m3), diameter (m) and height (m) of the MgO-C refractories before oxidation; V, R and h are the volume (m3), diameter (m) and height (m) of the unoxidized area following oxidation.
rm =
m0 m × 100% m0
(2)
where rm is the mass loss rate (%); m0 and m are the mass (g) before and after oxidation, respectively.
Fig. 1. Fabrication process of MgO-C refractories. 2
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Fig. 2. Temperature schedule of the coking (a) and oxidation (b) test.
Fig. 3. Bulk density (a) and apparent porosity (b) of cured and coked MgO-C refractories.
3. Results and discussion
oxidation rates of the MgO-C refractories were obtained by Eq. (1). The weight loss rates of MgO-C refractories were obtained by measuring the masses of MgO-C refractories before and after oxidation (Eq. (2)). The weight loss rate and oxidation rate results are shown in Fig. 6. From the figure, the weight loss rates increased while the oxidation rates decreased with increasing carbon content. Meanwhile, the weight loss rates and oxidation rates of MgO-C refractories fired at 1600 °C were higher than those fired at 1400 °C when the carbon content was 8 wt%, 12 wt%, and 16 wt%. However, abnormal results appeared when the carbon content was 3 wt%, and the weight loss rates and oxidation rates were lower at 1600 °C (marked by the red ellipse in Fig. 6). Due to the carbon content not changing the phase composition of MgO-C refractory, only the fired MgO-C refractory with 12 wt% carbon content was analysed in this study. The XRD patterns of the outer shell and the inner core are shown in Fig. 7. As can be seen from the figure, the main phases of the inner core were MgO, carbon, and MgAl2O4. A trace AlN phase and CaMgSiO4 impurity phase were also detected. The main phases of the outer shell were the same as in the inner core except without the carbon phase, which means that the carbon in the outer shell was completely oxidized. Meanwhile, the phase intensity of the MgAl2O4 was obviously increased in the outer shell samples. Aluminium powder reacted with nitrogen in the air to form AlN at high temperature (Eq. (3)). Also, part of the aluminium powder reacted with oxygen to form Al2O3 (Eq. (4)). The Al2O3 phase could also have been formed by the reaction between Al4C3 and CO (Eqs. (5) and (6)). The resulting alumina and magnesia formed MgAl2O4 (Eq. (7)).
3.1. Physical properties Fig. 3 shows the bulk density and apparent porosity of the MgO-C refractories with different carbon contents before and after coking. The bulk density of the cured MgO-C refractories decreased with increasing carbon content, which was mainly due to the lighter density of carbon than that of MgO. Meanwhile, the increasing of carbon content decreased the pores between MgO particles. After coking, the bulk density was decreased and the apparent porosity was increased significantly with the decomposing of phonetic resin to CH4, H2O, CO, C2H6, and H2 etc [17,18]. Different to the cured samples, the apparent porosity of the coked samples increased with the increasing of carbon content. Another noteworthy point is that the higher density and lower porosity were obtained in the 1400 °C batch. Fig. 4 shows the force-displacement curves and cold crushing strengths of MgO-C refractories before and after coking. From Fig. 4a–c, the failure displacements were decreased with the increasing of carbon content at 200 °C and 1000 °C, while it became ruleless when the coking temperature reached 1400 °C. From Fig. 4d, the cold crushing strength of the cured and coked MgO-C refractories decreased as the carbon content increased. Compared with the cured MgO-C refractories, the cold crushing strength of the coked MgO-C refractories were greatly reduced, and the higher the carbon content, the greater the reduction. 3.2. Oxidation resistance and mechanism The longitudinal section views of the fired MgO-C refractories are shown in Fig. 5. It can be seen that the MgO-C refractories have a distinct oxidized outer shell and an unoxidized inner core. Volumes of the shells and the cores were calculated from the above figure, and the 3
2Al(l, s)+N2(g)→2AlN(s)
(3)
4Al(l, s)+3O2(g)→2Al2O3(s)
(4)
4Al(l, s)+3C(s)→Al4C3(s)
(5)
Al4C3(s)+6CO(g)→2Al2O3(s)+9C(s)
(6)
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Fig. 4. Force-displacement curves: (a) cured at 200 °C, (b) coked at 1000 °C, (c) coked at 1400 °C and (d) cold crushing strength of MgO-C refractories.
Al2O3(s)+MgO(s)→MgAl2O4(s)
(7)
When oxidized at 1400 °C, the carbon in the MgO-C refractory reacted with oxygen to form carbon monoxide (direct oxidation, Eq. (8)), which formed the porous structure and isolated the MgO particles from each other in the decarburized layer (Fig. 8a-b). However, when the temperature was exceeded 1400 °C, the reaction between carbon and MgO started, which is the so-called indirect oxidation of carbon (Eq. (9)) [19]. The oxidation behaviour of the MgO-C refractory was the combined result of the direct oxidation and indirect oxidation of carbon. The reaction between MgO and carbon generated magnesium vapor, which was outwardly diffused through the pores of the decarburized layer where it reacted with oxygen to form MgO (Eq. (10)). The regenerated MgO partially filled the pores of the decarburized layer and connected the original MgO particles to form the continuous layer (Fig. 8c-d), which hindered the entry of the oxygen and improved the oxidation resistance of the MgO-C refractory to a certain extent [20]. The effect of indirect oxidation was more pronounced with lower
Fig. 6. Effect of carbon content on the oxidation rate and weight loss rate of MgO-C refractory.
Fig. 5. Longitudinal section view of the fired MgO-C refractories: (a) 1400 °C, (b) 1600 °C. 4
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Fig. 9. Effect of carbon content on the cold crushing strengths of fired MgO-C refractories. Fig. 7. XRD patterns of the fired MgO-C refractories: (a) 1400 °C, shell, (b) 1600 °C, shell, (c) 1400 °C, core, (d) 1600 °C, core.
is clear from Fig. 9 that, the cold crushing strengths fired at 1600 °C were higher than those fired at 1400 °C, especially for the fired MgO-C refractory with 3 wt% carbon content, which dramatically increased from 22.8 MPa to 56.5 MPa when the temperature increased from 1400 °C to 1600 °C. Undoubtedly, denser decarburized layer provided a better protection for the MgO-C refractory from the eroding away by the molten steel and slag [23].
carbon content in the MgO-C refractory, due to the MgO continuous layer being more easily formed when fewer pores were generated by the oxidation of carbon. Therefore, the MgO-C refractory with 3 wt% carbon content fired at 1600 °C had a lower weight loss rate and oxidation rate than that of fired at 1400 °C, even though the higher temperature means a faster reaction rate and mass transfer rate in general [21]. When the carbon content reached 8 wt% or higher, a more porous structure was created after oxidation of carbon, and the regenerated MgO was not sufficient to form a MgO continuous layer. Therefore, the effect of indirect oxidation was not so obvious and the weight loss rate and oxidation rate of the MgO-C refractories fired at 1600 °C were higher than those fired at 1400 °C. These results are in a good agreement with results of others [22]. 2C(s)+O2(g)→2CO(g)
(8)
C(s)+MgO(s)→CO(g)+Mg(g)
(9)
2 Mg(g)+O2(g)→MgO(s)
3.3. Oxidation kinetics Fig. 10 shows the cross section views of the MgO-C refractories after firing at 1000 °C for different durations. It can clearly be seen from the figure that the boundary of the non-oxidized core and the oxidized exterior shell are evident in four different carbon contents of MgO-C refractories. Meanwhile, the thicknesses of the oxidized exterior shells increased with the prolonging of firing time and decreased with the increasing of carbon content. According to the shrinking core model [24], the oxidation of MgO-C refractory can be divided into five parts: (Ⅰ) external mass transport of oxygen from the air flow bulk to the outer face of the sample, (Ⅱ) diffusion of oxygen through the pores in the decarbonized layer to reach the reaction interface, (Ⅲ) chemical reaction between graphite and oxygen at the reaction interface to form carbon monoxide, (Ⅳ)
(10)
The cold crushing strengths of the fired MgO-C refractories were also influenced by the formation of the MgO continuous layer (Fig. 9). It
Fig. 8. SEM images of the oxidized shell of fired MgO-C refractories: (a) 1400 °C, 3 wt% C, (b) 1400 °C, 16 wt% C, (c) 1600 °C, 3 wt% C, (d) 1600 °C, 16 wt% C. 5
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Fig. 10. Cross section views of fired MgO-C refractories: (a) 3 wt% C, (b) 8 wt% C, (c) 12 wt% C, (d) 16 wt% C.
diffusion of carbon monoxide through the decarbonized layer to the outer face, (Ⅴ) mass transfer of carbon monoxide from the outer face into the surrounding atmosphere. The schematic diagram of the MgO-C refractory oxidation process is shown in Fig. 11. When the temperature exceeded 800 °C, the external mass transport process of gas (parts Ⅰ and Ⅳ) and the reaction rate at the interface (part Ⅴ) were fast. Therefore, the gas diffusion in the decarburized layer can be deemed a restrictive factor and the oxidation rate equation of carbon can be shown as Eqs. (11) and (12) [25].
xC = 1
t=
rt r0
2
2 C r0 [x C + (1 4bDeff Cb
(11)
x C)ln(1
x C)]
(12)
where, t is the oxidation time (min), xC is the fractional oxidation of carbon at time t (%), ρC is the molar density of carbon in the MgO-C refractory (mol·cm−3), r0 is the initial radius of the MgO-C refractory (cm), b refers to the stoichiometric coefficient of carbon in the oxidation reaction with 1 mol of O2, Deff refers to the effective diffusion coefficient through the decarburized layer (cm2·min−1), and Cb is the oxygen concentration near the surface of the sample (mol·cm−3). According to the thicknesses of the decarburized layers measured
Fig. 11. Schematic diagram of the MgO-C refractory oxidation process.
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Fig. 12. Relationship between xC+(1-xC)ln(1-xC) and oxidation time. Fig. 13. Effect of carbon content on the effective diffusion coefficient and full oxidation times of MgO-C refractories.
Table 2 Related parameters of the oxidation rate equation for MgO-C refractories.
ρC × 10−2 (mol ⋅cm−3) r0 (mm) Cb × 10−6 (mol ⋅cm−3) b (-)
C1
C2
C3
C4
0.69 25 1.985 1.39
1.82 25 1.985 1.73
2.72 25 1.985 1.81
3.51 25 1.985 1.86
Table 3 Effective diffusion coefficient for previous oxidation tests.
from Fig. 10, the relationship between xC+(1-xC)ln(1-xC) and oxidation time can be obtained and is shown in Fig. 12. It can be seen that the two terms exhibits a good linear relation in the four different carbon contents of MgO-C refractories. The straight lines were fitted according to the dots and the slops were obtained. According to the related parameters of the oxidation rate equation shown in Table 2 and the definition of slope (Eq. (13)), the effective diffusion coefficient can be obtained. Meanwhile, the value of xC tended to the number 1 when the MgO-C refractory tended to totally oxidization, and the value of full oxidation time (tf) equalled to the reciprocal of the slope value (Eq. (14)).
4bDeff Cb S= 2 Cr0
tf =
2 1 C r0 = S 4bDeff Cb
Carbon content (wt %)
Temperature (oC)
Deff (cm2·min−1)
Reference number (-)
15 5 10 15 20 7 14 20
1000 1300
15.4 ~12 ~21 ~27 ~32 31.8 22.2 35.12 46.19 66.47 10.69–13.25 13.68–17.33
[14] [24]
14.3
1000 1000 1100 1200 950 1100
[26] [27] [28]
coefficients increased while the decarburized layer thicknesses decreased with the increasing of carbon content. This seemingly paradoxical result can be explained as follows. For a given decarburized layer thickness, the ratio of oxidation time (t) between x wt% carbon content (x is 3, 8, 12, 16) and 3 wt% carbon content of MgO-C refractory can be expressed as Eq. (16) [25]. The t was mainly related to the molar density of carbon (ρC) and effective diffusion coefficients (Deff) in the MgO-C refractories. Variations of ρC(x wt%)/ρC(3 wt%) and
(13) (14)
where, S is the slope of the straight line fitted by the experimental dots, tf is the full oxidation time (min). The results of effective diffusion coefficient and full oxidation time versus carbon content are shown in Fig. 13. It can be seen that the effective diffusion coefficient increased with the increasing of carbon content, which was mainly because the porosity of the decarburized layer was increased with the oxidation of carbon. This effect was partially impeded due to the reducing effect of carbon content on initial porosity of the MgO-C refractories (as seen in Fig. 3b). Meanwhile, the values of effective diffusion coefficients obtained in this study were much lower than that of others’ results (Table 3), which can be explained by the following: (1) 2 wt% of aluminium powder was added as antioxidant. The effect of aluminium powder on oxidation resistance of MgO-C refractory was not only on the preferential oxidization of aluminium over carbon, but also expansion in the specific volume accompanied by the formation of MgAl2O4 phase [29]. The expanded volume resulted in partial filling of the micro-pores of the decarburized layer and decreased inter diffusion rate of the gases through the solid and porous regions. (2) The hearth of the muffle furnace was a relatively closed space, where the mobility of the air around the MgO-C refractory was limited. As can be seen from Figs. 13 and Fig. 10, the effective diffusion
Fig. 14. Molar density ratio, effective diffusion coefficient ratio and full oxidation time ratio versus carbon content. 7
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Deff(x wt%)/Deff(3 wt%) with carbon content are shown in Fig. 14. It can be seen that both ratios increased with the increasing of carbon content. However, the increasing rate of the molar density ratio was higher than that of the effective diffusion coefficient ratio. Therefore, the oxidation time ratio increased with the increasing of carbon content. In other word, to form the same thickness of decarburized layer, longer oxidation time was required for MgO-C refractories with higher carbon content. For example, when the MgO-C refractory full oxidized (decarburized layer thickness of 25 mm), the full oxidation time was increased from 1339 min to 4566 min when the carbon content increased from 3 wt% to 16 wt% (Fig. 13). The increasing rate of the full oxidation time ratio was fell in between the increasing rate of the molar density ratio and the effective diffusion coefficient ratio (Fig. 14).
t (x wt%) C (x wt%)/ C (3 wt%) = t (3 wt%) Deff (x wt%)/ Deff (3 wt%)
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4. Conclusions Cured and coked MgO-C refractories were studied with respect to bulk density, apparent porosity, and cold crushing strength. The oxidation resistance and mechanism of MgO-C refractories were investigated via phase and microstructure evaluation. Also, the shrinking core model was applied in the oxidation kinetics of MgO-C refractories with different carbon content. Based on the above results, the following conclusions are drawn: (1) Bulk density, apparent porosity, and cold crushing strength of the cured MgO-C refractories decreased with the increasing of carbon content. Compared with the cured samples, the bulk density and cold crushing strength greatly reduced, and the apparent porosity significantly increased after curing at 1000 °C and 1400 °C. (2) Mg vapor was generated by the reaction between MgO and carbon at 1600 °C, which re-oxidized to MgO in the outward diffusion process of the decarburized layer. The regenerated MgO resulted in apparent porosity reduction and pore-filling effects, which caused diminution of the internal diffusion of gases in the decarburized layer. Finally, the oxidation resistance of MgO-C refractory with 3 wt% carbon content fired at 1600 °C was higher than that fired 1400 °C. (3) Effective diffusion coefficient increased while the thickness of the decarburized layer decreased with the increasing of carbon content. The seemingly paradoxical result was mainly due to the effect of the molar density of carbon prevailing over the effective diffusion coefficient in the oxidation process. As a result, longer oxidation time was required for full oxidation of MgO-C refractories with higher carbon content. 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. Acknowledgements The authors would like to express their gratitude for the financial support from the National Natural Science Foundation of China (Grant No. 51974074, 51874083 and 51404056) and the Fundamental
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