Dissolution behavior of a novel Al2O3-SiC-SiO2-C composite refractory in blast furnace slag

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Dissolution behavior of a novel Al2O3-SiC-SiO2-C composite refractory in blast furnace slag ⁎

Haibin Zuoa, , Cong Wangb, Yingli Liua a b

State Key Laboratory of Advanced Metallurgy, University of Science and Technology Beijing, Beijing 100083, China Patent Examination Cooperation Tianjin Center of the Patent Office, SIPO, Tianjin 300304, China

A R T I C L E I N F O

A BS T RAC T

Keywords: Blast furnace Refractories Corrosion Dissolution

Al2O3-SiC-SiO2-C composite refractories are interesting potential blast furnace hearth lining materials that feature several advantageous properties. In this study, the corrosion resistance of a novel Al2O3-SiC-SiO2-C composite refractory to blast furnace slag was investigated by adopting a rotating immersion method (25 r/min) at 1450–1550 °C and comparing it against a conventional corundum-based refractory at 1550 °C as a benchmark. The results showed that the apparent activation energy of Al2O3-SiC-SiO2-C composite refractory over the dissolution process in the slag is 150.4 kJ/mol. Dissolution of the Al2O3 and 3Al2O3·2SiO2 phases appeared to be the main cause of Al2O3-SiC-SiO2-C composite refractory corrosion. High-melting-point compounds in the slag layer formed a protective layer which mitigated the corrosion. The novel Al2O3-SiCSiO2-C composite refractory is better suited to blast furnace hearth lining than the conventional corundumbased refractory, because the carbon phase and SiC phase in the material are not readily wetted by the blast furnace slag and therefore are more resistant to slag penetration. Higher melting point phases also may crystallize on the hot face of the hearth lining due to the high thermal conductivity of the Al2O3-SiC-SiO2-C composite refractory, promoting a more stable protective layer.

1. Introduction Over the past decade, many steel companies have suffered severe safety accidents due to blast furnace hearth issues such as burning through of the hearth sidewall or abnormal temperature increase in the carbon brick. To this effect, improving existing hearth refractories has become an increasingly urgent endeavor [1,2]. Oxidation at high temperatures, corrosion by smelting slag and iron, and heat stress damage are destructive to the hearth line. Traditional hearth materials such as carbon brick or corundum brick (ceramic cup) are inherently unsafe unless combined with other, more advantageous materials to construct hearth structures. Composite structures can be utilized to design a safer hearth by minimizing erosion; once the protective layer brick (typically ceramic cup) is damaged, however, the remaining refractory is exposed to smelting slag. Alumina-carbon refractories are often utilized as blast furnace hearth lining due to their favorable thermal conductivity and thermal shock resistance [3–5]. Molten slag corrosion is an important threat to the hearth lining, so slag resistance is one of the most important properties of hearth refractories. Carbon phase has high thermal conductivity and is not wetted by slag [6,7], so it is often added to

⁎

refractories to increase their corrosion resistance and thermal conductivity [8]. The presence of carbon phase also promotes oxidation corrosion in the refractory due to oxygen in the atmosphere [9,10]. SiC phase (typically via the addition of SiC or Si) is also often introduced into alumina-carbon refractories in order to enhance their oxidation resistance and decrease their pore size [11,12]. Al2O3-SiC-SiO2-C composite refractories have, accordingly, gradually become preferable materials for blast furnace heath lining as they possesses excellent thermal conductivity and slag corrosion resistance at high temperatures. A fundamental understanding of refractory performance at high temperatures is the best approach to decreasing energy consumption, reducing downtime, and increasing the campaign life of blast furnaces. That being said, there have been very few studies related to the corrosion mechanisms of Al2O3-SiC-SiO2-C composite refractories by molten slag, especially compared to the extant research on MgO-C refractories [13–17] and spinel-containing refractories [18–21]. There have been valuable contributions to the literature, however: Lee et al. [22], for example, studied the dissolution kinetics of dense alumina in calcium aluminosilicate-based melts at 1560–1590 °C to find that dissolution rate is most likely controlled by mass transfer in the slag

Corresponding author. E-mail address: [email protected] (H. Zuo).

http://dx.doi.org/10.1016/j.ceramint.2017.02.138 Received 22 October 2016; Received in revised form 6 February 2017; Accepted 26 February 2017 0272-8842/ © 2017 Elsevier Ltd and Techna Group S.r.l. All rights reserved.

Please cite this article as: Zuo, H., Ceramics International (2017), http://dx.doi.org/10.1016/j.ceramint.2017.02.138

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phase; Hong et al. [23] investigated the interactions of Al2O3-SiC-SiO2C composite refractory with slag and found that low-melting-point compounds formed on the material surface representing significant damage to the refractory. Yang et al. [24] used fractal theory to study the corrosion resistance of unfired Al2O3-SiC/β-Sialon/Ti(C, N)-C refractory, and found that β-Sialon and SiC powders together improve slag corrosion resistance performance. Justus et al. [25] investigated the corrosion mechanism of Al2O3-SiC-C-MgAl2O4 refractory to find that slag containing large quantities of calcium aluminum silicate interact with the refractory microstructure to cause corrosion. Chan et al. [26] explored the influence of additives on the slag resistance of Al2O3-SiC-SiO2-C refractory bond phases under a reducing atmosphere, and found that carbon content dominates the resistance to molten slag penetration. These studies mainly focused on the slag corrosion resistance of refractories under static conditions despite the fact that the service conditions of most refractories are dynamic. Slag corrosion resistance properties under accurate, dynamic production conditions thus merit further research. In this study, the slag corrosion resistance of a novel Al2O3-SiCSiO2-C composite refractory was investigated via rotating immersion method. The kinetics of the dissolution process were quantified by measuring the decrease in sample radius as a function of corrosion exposure duration at various temperatures. We also tested a corundum-based refractory at 1550 °C as a benchmark for comparison against the proposed refractory. The interactions between the refractory and slag were also investigated to elucidate the corrosion mechanism. Our primary goal is to develop a workable understanding of the corrosion behavior and kinetics of the Al2O3-SiC-SiO2-C composite refractory in contact with molten slag, and ultimately, to expand the campaign lives of existing blast furnaces.

Table 2 Physical and chemical properties of Al2O3-SiC-SiO2-C composite refractory and corundum-based refractory. Item

Unit

Al2O3-SiC-SiO2C composite refractory

Corundumbased refractory

Bulk density Apparent porosity Oxidation ratio (Gas) Permeability Resistance to molten iron infiltration Average pore diameter Pore volume ( < 1 µm)

g/cm3 % % mDa %

2.98 10.9 110 0.63 0.12–0.62

3 15 110 1.0 1.0

µm %

0.238 80.44

1.0 70

Thermal conductivity

W/(m °C)

17.34 16.21 14.27 13.78

5.91 – 4.53 4.09

20 °C 300 °C 600 °C 800 °C

2. Experimental procedure 2.1. Experimental materials and methods The samples of refractories and all raw materials were provided by Henan Winna Industrial Group Co., Ltd. The chemical compositions of the Al2O3-SiC-SiO2-C composite and corundum-based refractory are presented in Table 1. The Al2O3-SiC-SiO2-C composite was made commercially available only recently; it is a novel refractory material that has been successfully used in blast furnace hearth applications exhibiting high refractoriness, high thermal shock resistance, and low wettability by molten slag. The physical and chemical properties of the Al2O3-SiC-SiO2-C composite refractory and corundum-based refractory are summarized in Table 2. A schematic diagram of the experimental apparatus used in the slag resistance experiment is shown in Fig. 1. The Al2O3-SiC-SiO2-C composite refractories were machined into cylindrical samples (external diameter=18 mm; height=20 mm) with a through-bore (diameter=4 mm). The samples were then dried at 110 °C for 24 h to remove any water. Screw threads (length=50 mm) were machined at one end of the Mo bars (diameter=4 mm; height=500 mm), which were then fixed to each cylindrical sample at both ends. The Mo bar was expected to provide enough rigidity at high temperature to keep the sample stable and to push the sample bar into the molten slag. Blast furnace slag (125 g) was placed into a Mo crucible (inside diameter=38 mm; height=60 mm) that was held inside a graphite crucible (inside

Fig. 1. Schematic of experimental set-up.

diameter=39; height=80 mm) to prevent the slag from splashing. High-purity (99.9 vol%) argon gas was purged through the reaction tube at a flow rate of 2 L/min during the experiment to prevent oxidation of the sample or graphite crucible. Once the desired temperature was reached, the Mo crucible containing the test slag was placed in the constant temperature zone in the experimental furnace. The furnace temperature was stabilized for 30 min to ensure that the test slag melted in the Mo crucible, then the Mo bar connected to the fixed sample was placed into the molten slag and rotated at a constant speed of 25 r/min. The outside surface of the cylinder was exposed to the molten slag while the sample was rotated to simulate molten slag flow in the blast furnace hearth lining. After the required exposure time, the sample was raised 5 cm above the Mo crucible and rotated at a high speed (100 r/min) for 4 min to remove the molten slag adhering to the sample surface. The sample was then quenched in air and prepared for further analysis. As mentioned above, we also tested a corundum-based refractory under the same experimental methodology for the sake of comparison against the proposed Al2O3-SiC-SiO2-C composite refractory. Table 3 shows the detailed experimental scheme.

Table 1 Chemical analysis results of the Al2O3-SiC-SiO2-C composite refractory and corundumbased refractory, wt%. Item

Al2O3

C

SiO2

SiC

Others

Al2O3-SiC-SiO2-C composite refractory Corundum-based refractory

73.1 83.0

10.2 3.6

8.2 9.7

6.0 3.6

2.5 –

2.2. Characterization The sample diameter was measured at several points with a slide caliper before and after the experiment. The mean decrease in the radius was estimated based on this data and used to calculate the mass transfer coefficient for the dissolution process. After these measure2

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ments were complete, smaller samples were cut from the corroded area of the full test sample and carefully prepared for phase characterization and microstructure analysis at the slag/refractory interface. Test slag, blank sample, and test sample phase characterizations were performed via X-ray diffraction (XRD, UltimaIV, Cu Kα radiation at 40 kV) (Rigaku, Tokyo, Japan). The X-ray fluorescence (XRF) method was used to determine the test slag content. The viscosity of the blast furnace slag was measured with a comprehensive properties measurement instrument (RTW-10, Northeastern University, Shenyang, China). A scanning electron microscope (SEM) (FEI Quanta 250, FEI, Hillsboro, USA) equipped with energy dispersive X-ray spectroscopy (EDS) was used to measure the morphology and elemental distribution of the blank and test samples.

silicon powder existed in the raw materials), because the silicon had far lower reaction starting temperature with carbon to form SiC compared to the reduction of silica. Volume-expansive effect due to in situ SiC formation generated an abundance of pores during roasting that were subsequently filled and sealed off, decreasing the porosity and average pore diameter of the product and increasing the proportion of pores less than 1 µm in size. The presence of SiC phase also enhanced the oxidation resistance of the material due to the formation of SiO2 during oxidation. The Al2O3 phase in the refractory directly originated from the fused alumina (Al2O3% > 98.5%) in our initial ingredients which constituted the refractory matrix. We readily deduced the 3Al2O3·2SiO2 as a product of the reaction between Al2O3 and SiO2. The presence of Al2O3 and 3Al2O3·2SiO2 increased the material's chemical stability, including its excellent resistance to molten iron. Fig. 3 shows the SEM micrograph and EDS spectra of the Al2O3SiC-SiO2-C composite refractory, where the material structure is quite dense and without apparent pores on the surface. According to the EDS spectra, the carbon phases were diffusely distributed in the Al2O3 matrix containing some 3Al2O3·2SiO2. These carbon phases appeared to be surrounded by Al2O3 or 3Al2O3·2SiO2, which is the main reason the Al2O3-SiC-SiO2-C composite refractory showed such excellent oxidation resistance compared to the carbon bricks.

3. Results and discussion

3.2. Slag characterization

3.1. Phase compositions and microstructure of Al2O3-SiC-SiO2-C composite refractory

Our chemical analysis of the experimental blast furnace slag determined by XRF is shown in Table 4. The basicity (CaO/SiO2) index of the slag was 1.08. Our XRD analysis of the blast furnace slag (crystallized during quenching) is shown in Fig. 4. Most phases in the slag were in an amorphous state and most of the crystalline phase was akermanite-gehlenite that had formed during the crystallization process. Viscosity is one of the most important physical properties of blast furnace slag and significantly affects the extent of refractory corrosion: Decrease in slag viscosity can aggravate refractory corrosion due to increase in the flow velocity of the slag on the refractory surface. Fig. 5 shows the viscosity of the experimental slag at various temperatures. The viscosity increased as temperature decreased, suggesting that the refractory was easily corroded at elevated temperatures. We calculated the slag melting temperature at 1378 °C with a corresponding viscosity of 0.965 Pa s.

Table 3 Experimental scheme. Sample

Temperature,°C

Corrosion time, min

Al2O3-SiC-SiO2-C composite refractory

1450 1500 1550 1550

30, 60, 90, 120 30, 60, 90,120 30, 60, 90, 120 120

Corundum-based refractory (contrast test)

The phase compositions of Al2O3-SiC-SiO2-C composite refractory were determined by XRD at 25 °C. The XRD pattern, as shown in Fig. 2, contained C, Al2O3, SiC, and 3Al2O3·2SiO2 (Mullite). The carbon phase was majorly derived from natural flake graphite powder (particle size < 149 µm and C% > 95%) which was added into mixture during batching to improve the thermal conductivity of refractory. The presence of natural flake graphite is also known to increase the corrosion resistance of refractories, contributing to their non-wetting characteristics towards molten slag. The SiC phase in the refractory was mostly generated by the reaction during roasting between silicon powder and carbon black (particle size < 1 µm and C% > 98%) and graphite, which were added in the initial batching stage. We could not completely neglect the possibility that SiC was generated by the reactions between silica or mullite and carbon, as observed by Luz and Pandolfelli in a similar experiment; there were certainly many factors influencing SiC phase generation during the preparation process [27]. Roasting temperature was lower than 1700 °C in our manufacturing process, so we do not believe the reduction of silica was a significant source of SiC (especially when 7000

β α

6000

Intensity, a.u.

The decreases in radius of Al2O3-SiC-SiO2-C composite refractories after being corroded by the slag between 1450 and 1550 °C are plotted as a function of immersion time in Fig. 6. We found that dissolution depth increased as corrosion exposure time increased. Accordingly, we calculated the dissolution rate of the samples in molten slag at different temperatures to be 0.00617 mm/min (1450 °C), 0.00453 mm/min (1500 °C), and 0.00807 mm/min (1550 °C). Many factors can influence corrosion rate, but dissolution is of particular importance [8]. Theoretically, the dissolution rate can be controlled by either chemical reactions or mass transport of reaction species through the molten slag at the slag-refractory reaction interface. The slope in Fig. 6 indicates a linear relationship between dissolution depth and corrosion exposure time, which further indicates that the dissolution rate was controlled via mass transport of the reaction species [28]. The decrease in sample radius (−dr/dt) is also likely related to the dissolution rate of the refractory into the molten slag, as expressed in Eq. (1) based on the Arrhenius equation [29]:

α—Al2O3

α

β—3Al2O3·2SiO2

δ β

5000

3.3. Corrosion test

χ—SiC δ—C

α

α

4000 3000

α

2000 1000

β

β

0 20

β

χ

β ββ δ 40

α

α β α

δ β β χ αβ β αβ δ χα α αααα α β β ββ β 60

80

100

J=−D

2θ, degree

E Δc dr = ρr ( − ) = Aexp( − c ) dt RT δ

(1)

where J is the mass flux for the dissolution (g/(cm s)), D is the 2

Fig. 2. XRD pattern of Al2O3-SiC-SiO2-C composite.

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Fig. 3. SEM micrograph and EDS spectra of the Al2O3-SiC-SiO2-C composite refractory. Table 4 Chemical composition of experimental blast furnace slag, wt%. CaO

SiO2

Al2O3

MgO

TiO2

Total Fe

S

MnO

Others

35.05

32.47

18.04

7.65

1.95

0.82

0.44

0.87

2.71

1600

θ:Akermanite-gehlenite

θ

1400

Intensity, a.u.

1200 1000 800

Fig. 5. Relation between the slag viscosity and temperature.

600 400 200

θ θθ θ θ

θ

θ θ θ θ θ θθ θ θ θθθ θ θθ θ θ θ θθ

diffusivity of the sample (cm2/s), Δc is the difference in concentration between the interface and the equilibrium slag bulk (g/cm2), δ is the boundary layer thickness (cm), ρr is the density of the sample (g/cm3), A is the pre-exponential factor, and R is the universal gas constant, 8.314 J/(mol K). Because the decrease in sample radius (−dr/dt) can be used to characterize the dissolution rate at different temperatures, Eq. (1) can be transformed as follows:

θ

0 20

40

60

80

100

2θ, degree Fig. 4. XRD pattern of the blast furnace slag.

4

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Decrease in radius, mm

1.0

0.8

refractory. We used the SEM photo and EDS mapping results of the sample at 1550 K for 120 min immersion time to directly analyze the corrosion interface (Fig. 9). The distributions of Al2O3 and 3Al2O3·2SiO2 were ascertained from this figure, as well. Al and O were detected in the matrix but not Si, which marks the presence of Al2O3. In addition, Si, Al, and O were detected simultaneously, which marks the 3Al2O3·2SiO2 phase. Fig. 9 also shows where Ca and Mg elements were distributed along the inside surface of the sample as a result of the slag penetrating into the refractory. Some small Al2O3 particles were present in the slag close to the boundary of undissolved Al2O3 particles, indicating that the Al2O3 particles had gradually dissolved in the molten slag during the corrosion process. The corrosion mechanisms of Al2O3 and 3Al2O3· 2SiO2 in our sample were penetration, dissolution, refractory-slag reactions leading to a transformation of the refractory microstructure, and finally, erosion of the refractory. However, the viscosity of the slag adjacent to the sample increases as more and more Al2O3 constantly dissolves into the slag [30], which is beneficial in terms of reducing the erosion damage by decreasing the slag flow rate on the sample surface. The regions in which carbon phases were present were less corroded than the regions without carbon phase (Fig. 9). As mentioned above, carbon phase is not readily wetted by the molten slag, thus the resistance to slag penetration increased in the areas containing carbon phase. SiC also has excellent slag corrosion resistance [31], which further contributed to the slag corrosion resistance of the Al2O3-SiCSiO2-C composite refractory. Compared to carbon and SiC phases, the mullite phase actually damages the corrosion resistance of the refractory; the CaO and MgO in the slag were the main cause of corrosion to the mullite phase. CaO and MgO penetration promoted the decomposition of the mullite phase. The decomposition product SiO2 and Al2O3 reacted with CaO and MgO, then generated new phases such as Ca2Al2SiO7, Ca2MgSi2O7, and MgAl2O4. These new phases broke away from the refractories as the slag was constantly flushed, leading to the destruction and corrosion of the refractories. In order to analyze the phase composition of the slag layer after corrosion, we carefully collected the mixtures attached to the sample after rapid cooling and analyzed them via XRD as shown in Fig. 10. The slag layer appeared to be a mixture of akermanite (2CaO·MgO·2SiO2)gehlenite (2CaO·Al2O3·SiO2), carbon (C), corundum (Al2O3), and spinel (MgO·Al2O3). Unreacted Al2O3 phase and C phase were detected in the slag layer due to the dissolution of the sample; 2CaO·MgO·2SiO2 phase, 2CaO·Al2O3·SiO2 phase, and MgO·Al2O3 phase were also observed in the slag layer having been generated by various reactions (Eqs. (3)–(5)) with melting points of 1450 °C, 1593 °C, and 2135 °C respectively.

1450 oC 1500 oC 1550 oC

0.6

0.4

0.2

0.0 30

60

90

120

Corrosion exposure time, min Fig. 6. Relations between the decrease in radius of the samples and corrosion time at different temperatures.

ln( −

E dr A ) = − c + ln dt RT ρr

(2)

This equation was used to calculate the activation energy Ec. The rate constant for the dissolution process at the surface of the Al2O3-SiC-SiO2-C composite refractory is affected by the temperature of the blast furnace hearth. This dependence can be described according to Eq. (2), as can the fit of dissolution rates with various temperatures; the best fit line we drew according to this equation is shown in Fig. 7. The activation energy Ec for the dissolution rate of the Al2O3-SiC-SiO2-C composite refractory into the blast furnace slag was estimated to be 150.4 kJ/mol – much greater than the 26.4 kJ/mol obtained from the dissolution test of a MgO-C brick in a CaO–Al2O3– SiO2–MgO slag system by Kasimagwa et al. [8]. Fig. 8 shows the microstructure of the reaction interface between the refractory and slag at 1500 °C for 2 h. The concentration profiles of the main elements (along the yellow line) were analyzed by EDS to find that the penetration of slag into the refractory microstructure mainly took place at the Al2O3 or 3Al2O3·2SiO2 particle boundaries. In general, calcium in the blast furnace slag is the main element penetrating into the interior structure of the matrix. Consequently, the calcium content in both the molten slag and matrix can be measured to estimate the extent to which the matrix is corroded by the melting slag. According to the line scan shown in Fig. 8, the calcium content in our sample decreased continuously from the slag layer to the matrix. The concentrations of magnesium and silicon decreased sharply from the slag layer to the matrix because these elements existed in the slag. By contrast, the aluminum concentration increased sharply from the slag layer to the matrix due to the inherently high aluminum content in the

2CaO+MgO+2SiO2=2CaO∙MgO∙2SiO2

(3)

2CaO+Al2O3 + SiO2=2CaO∙Al2O3∙SiO2

(4)

MgO+Al2O3 = MgO∙Al2O3

(5)

The reactions described by Eqs. (6)–(9) may also have occurred during the corrosion process. These products were not present in the XRD patterns, however, as they generated relatively scant amounts of oxidation products.

2CaO+SiO2 = 2CaO∙SiO2

(6)

CaO + 2Al2O3=CaO∙2Al2O3

(7)

CaO+Al2O3=CaO∙Al2O3

(8)

CaO+Al2O3+2SiO2=CaO∙Al2O3∙2SiO2

(9)

The melting points of the 2CaO·SiO2, CaO·2Al2O3, CaO·Al2O3, and CaO·Al2O3·2SiO2 reaction products are 2130 °C, 1600 ℃, 1780 °C, and 1557 °C, respectively. In effect, the melting points of actual products and possible products in the slag layer are higher than or similar to the temperature in the blast furnace hearth. The Al2O3-SiC-SiO2-C com-

Fig. 7. Effect of temperature on rate constant for dissolution process of the samples.

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Fig. 8. SEM photo and line scanning result of the sample contacting with slag at 1500 °C for 120 min. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)

Fig. 9. SEM photo and EDS mapping results of the sample contacting with slag at 1550 °C for 120 min.

high-melting-point phases leads to the formation of a protective layer which prevents further corrosion. There are many factors affecting the generation of new phases associated with impurities, mainly TiO2, Fe2O3, K2O, and Na2O. An array of factors further affect the quantity and nature of any surrounding phase. These impurities are common in real-world blast furnace slag originating from the ironmaking materials; refractory corrosion and dissolution does not alter the contents of these impurities near the corroding surface to any significant extent. Although some low-melting-point detrimental phases did appear due to impurities during corrosion, the increase in Al2O3 content on the corroding surface played a dominant role in crystallization as-evidenced by the XRD patterns (Fig. 10). We also observed a few MgAl2O4 spinel phases in the slag layer attached to the corroding surface during the corrosion process. These phases may have been generated from the reactions between MgO in the slag and Al2O3 in the refractories or slag. The in-situ MgAl2O4 spinel formation on the refractory surface caused structural damage to the refractories due to its volume-expansive effect, subsequently accelerating the pulverization and breaking from the corroding surface during corrosion [32]. The amount of MgAl2O4 spinel generated during our tests was minimal, so its destructive effect was rather limited. We compared the proposed Al2O3-SiC-SiO2-C composite refractory

posite refractory possesses high thermal conductivity due to the addition of C phase, which can decrease the hot face temperature of the hearth refractory coupled with the effect of cooling water on the outside furnace wall. Once the hot face temperature drops below the melting point of the phases described above, the crystallization of the

3000

Intensity, a.u.

θ—Ca2MgSi2O7-Ca2Al2SiO7

δ

4000

φ—MgAl2O4 δ—C α—Al2O3

φ θ

2000

1000

0

θφ δ θ θ θ φα θ θφ δ α θ θ α θ θαθ θθφθ α θα θ θ θ αθ θ φ θ θθθ 20

40

60

α θ 80

2θ, degree Fig. 10. XRD pattern of the slag layer adhering to the sample surface after corrosion at 1550 °C for 120 min.

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Fig. 11. SEM photo and EDS mapping results of a traditional corundum-based refractory contacting with slag at 1550 °C for 120 min.

slag adjacent to the sample, which mitigated erosion damage by decreasing the slag flow rate on the sample surface. High melting points were observed in the slag layer (2CaO·MgO·2SiO2, 2CaO·Al2O3· SiO2, MgO·Al2O3) which formed a protective layer that prevented further corrosion. The proposed Al2O3-SiC-SiO2-C composite refractory is better suited to blast furnace hearth lining than the conventional corundum-based refractory because 1) the carbon and SiC phases in the material are not readily wetted by the molten slag, thus enhancing the resistance to slag penetration; 2) the former has higher thermal conductivity due to the presence of C phase while the hot face temperature of the refractory can be decreased significantly through cooling water; 3) more high-melting-point phases may be crystallized on the proposed refractory surface due to low hot-surface temperature, which promotes a more stable protective layer.

to the conventional blast furnace hearth refractory by also testing a corundum-based refractory under the same experimental methodology described in Section 2. Table 1 shows where the effects of SiC and C were negligible due to their low quantities in the material. The SEM photo and EDS mapping results of the corundum-based refractory after corrosion are shown in Fig. 11. In the SEM photo, the corrosion boundary appears to fall between the slag layer and the matrix. Some undissolved Al2O3 particles are also apparently distributed in the slag layer close to the matrix, and some pores exist in the slag layer which suggest that it was not stable at high temperature. The EDS mapping results indicate that the distributions of Ca and Si elements outline the corrosion zone (consistent with the SEM results). The Al content in the matrix exceeded that in the slag layer due to the corundum-based refractory containing a large amount of Al2O3. There was O element present in both the slag layer and matrix because most phases in the slag and the material are oxide masses. Taken together, our results indicate that the corrosion resistance of Al2O3-SiC-SiO2-C composite refractory is better than that of corundum-based refractory due to the addition of a C phase which is not wetted by molten slag. Appropriate amounts of carbon and silicon carbide can balance the thermal conductivity and corrosion resistance of blast furnace refractories. Excessive amounts of either lead to a decline in the material's performance (e.g., mechanical properties, oxidation resistance), however.

Acknowledgements This study was conducted with financial support from the National Natural Science Foundation of China (No. 51574023). References [1] F.M. Zhang, Design and operation control for long campaign life of blast furnaces, J. Iron Steel Res. Int. 20 (2013) 53–60. [2] Z.J. Liu, J.L. Zhang, H.B. Zuo, T.J. Yang, Recent progress on long service life design of Chinese blast furnace hearth, ISIJ Int. 52 (2012) 1713–1723. [3] Y. Xu, S. Sang, Y. Li, B. Ren, et al., Pore structure, permeability, and alkali attack resistance of Al2O3-C refractories, Metall. Mater. Trans. A 45 (2014) 2885–2893. [4] A.P. Luz, M.M. Miglioli, T.M. Souza, et al., Effect of Al4SiC4 on the Al2O3-SiC-SiO2C refractory castables performance, Ceram. Int. 38 (2012) 3791–3800. [5] S. Zhang, W.E. Lee, Carbon containing castables: current status and future prospects, Br. Ceram. Trans. 101 (2002) 1–8. [6] N. Prompt, E. Ouedraogo, High temperature mechanical characterisation of an alumina refractory concrete for blast furnace main trough: Part I. General context, J. Eur. Ceram. Soc. 28 (2008) 2859–2865. [7] E. Ouedraogo, N. Prompt, High-temperature mechanical characterisation of an alumina refractory concrete for blast furnace main trough: Part II. Material behavior, J. Eur. Ceram. Soc. 28 (2008) 2867–2875. [8] I. Kasimagwa, V. Brabie, P.G. Jönsson, Slag corrosion of MgO-C refractories during secondary steel refining, Ironmak. Steelmak. 41 (2014) 121–131. [9] H. Zuo, C. Wang, J. Zhang, et al., Oxidation behavior and kinetics of Al2O3-SiCSiO2-C composite in air, Ceram. Int. 41 (2015) 9093–9100. [10] H. Zuo, C. Wang, J. Zhang, et al., Comparison of oxidation behaviors of novel carbon composite brick with traditional carbon brick, Ceram. Int. 41 (2015) 7929–7936. [11] T. Goto, H. Homma, High-temperature active/passive oxidation and bubble formation of CVD SiC in O2 and CO2 atmospheres, J. Eur. Ceram. Soc. 22 (2002)

4. Conclusions The corrosion resistance of a novel Al2O3-SiC-SiO2-C composite refractory in blast furnace slag was investigated in this study via rotating immersion method (25 r/min) under dynamic conditions at 1450–1550 °C. The dissolution rates of several samples in the slag at various temperatures were 0.00617 mm/min (1450 °C), 0.00453 mm/ min (1500 °C), and 0.00807 mm/min (1550 °C). The corresponding apparent activation energy obtained for the dissolution process was 150.4 kJ/mol, which may be an important reference for future researchers. A slag layer adhered to the sample surface after the corrosion experiment. Penetration of the slag into the refractory microstructure was observed mainly at the boundary of Al2O3 or 3Al2O3·2SiO2 particles. The dissolution of Al2O3 and 3Al2O3·2SiO2 was found to be the main corrosion mechanism of the Al2O3-SiC-SiO2-C composite refractory, though Al2O3 dissolution also increased the viscosity of the 7

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