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Ceramics International 40 (2014) 9763–9773www.elsevier.com/locate/ceramint

Preparation and blast furnace slag corrosion behavior of SiC–Sialon–ZrNfree-fired refractories

Juntong Huang, Zhaohui Huangn, Yan'gai Liu, Minghao Fang, Haitao Liu, Xiaowei Cao,Xiaochao Li, Mengyan Yin, Ruilong Wen, Hao Tang

School of Materials Science and Technology, China University of Geosciences (Beijing), Beijing 100083, P. R. China

Received 15 January 2014; received in revised form 15 February 2014; accepted 16 February 2014Available online 24 February 2014

Abstract

Sialon–ZrN powders were synthesized from low-grade bauxite using a zirconite additive by carbothermal reduction nitridation (CRN). The as-synthesized Sialon–ZrN powders were subsequently used as SiC-based free-fired refractories. Their phase compositions and microstructures ofthe powders were determined using X-ray diffraction and scanning electron microscopy. The physical properties and blast furnace (BF) slagresistance of the SiC–Sialon–ZrN free-fired refractories were also studied. The results show that the low-grade bauxite and zirconite powderswere transformed to β-Sialon and ZrN by the CRN process at 1600 1C. The phenolic resin provided a strongly bonded interface between theSialon–ZrN matrix and SiC particles and thus enhanced their combined strength after drying at 150 1C. The strength increased as the temperaturewas elevated from 1000 1C to 1500 1C. As the ZrN content increased, the slag erosion rate of the SiC–Sialon–ZrN free-fired refractories initiallydecreased and then increased after being heated at 1500 1C. The presence of ZrO2 in the slag (oxidized from ZrN) revealed that the ZrO2 did notreact with other oxides in the BF slag to form low melting point phases. This may play a crucial role in its BF slag erosion resistance.The corrosion mechanism of the BF slag towards the as-prepared SiC–Sialon–ZrN free-fired refractories was determined to be oxidation-erosion-dissolution-penetration.& 2014 Elsevier Ltd and Techna Group S.r.l. All rights reserved.

Keywords: Sialon–ZrN powders; Low-grade bauxite; Carbothermal reduction-nitridation; SiC-based free-fired refractories; Blast furnace slag resistance

1. Introduction

Sialon-bonded SiC composite refractories have better abra-sion, alkali, and thermal shock resistances than high-aluminaor corundum bricks. They also have better oxidation resistanceand higher strength than carbon products. Thus, they arewidely applied, especially as blast furnace (BF) refractories (acritical and dominating factor in the quality of iron smeltingand also the lifespan of blast furnace) [1–3].

As the matrix in Sialon–SiC composite refractories, Sialon hasmostly been prepared using Si, Al, Al2O3 and/or Si3N4 powdersas raw materials using an in situ process of nitridation reactionsintering [3,4]. The high cost of the raw materials seriously limits

10.1016/j.ceramint.2014.02.06314 Elsevier Ltd and Techna Group S.r.l. All rights reserved.

g author. Tel./fax: þ86 10 82322186.ss: huang118@cugb.edu.cn (Z. Huang).

its industrial applications as a custom refractory material. In1979, Lee et al. [5] prepared Sialon using natural kaolin bycarbothermal reduction-nitridation (CRN) reaction. Since then,this method has offered a new technical path to prepare Sialon-based materials economically by carbothermal conversion fromnatural products and waste materials containing Si and/or Alelements. Materials such as quartz [6], natural clay [7], coalgangue [8] and fly ash [9] could then be used instead ofrelatively expensive synthetic powders. Low-grade bauxite is anore with a lower content of Al2O3 content and occurs just abovethe bauxite layers in mines. Composed of SiO2 and Al2O3, itconstitutes an ideal and cheap raw material for synthesizingSialon [10]. China has rich bauxite resources, especially inYangquan, Shanxi province, and there is much more low-gradebauxite containing only very few impurities, e.g., iron and basicmetals. The low-grade ores are usually just piled up, not only

J. Huang et al. / Ceramics International 40 (2014) 9763–97739764

resulting in pollution of the environment, but also in wastage ofthe resources. Thus, it is necessary to increase the use of suchlow-grade bauxite in high value-added products and using it as araw material to synthesize Sialon makes good environmental andeconomic sense.

Previous studies have suggested that the slag and thermalshock resistances of Sialon–SiC composites, however, requiresome improvement. It is well known that zirconium nitride(ZrN) possesses excellent performances. These include highmelting point, high strength and toughness, and also good slagresistance and stability towards thermal shock [11]. ZrN canalso serve as an antioxidant. Even though it can be oxidized toZrO2 at high temperature, it is not easily wetted by metallic orBF slag. Therefore, we decided to introduce ZrN into Sialon–SiC composites, wherein the Sialon and ZrN would besynthesized from low-grade bauxite and zirconite using theCRN method.

Fabrication of Sialon–SiC based refractories has previouslyemployed in situ nitridation reaction sintering at high temperature.This method, however, suffers from high energy consumption. Atpresent, energy conservation and emission reduction are essentialto develop economic prosperity in harmony with environmentalconcerns. Therefore, for this work, we preferred to use a free-firedtechnique to prepare the refractory materials. First, the Sialon–ZrNcomposite powders were synthesized from low-grade bauxite andzirconite using the CRN method, and their phase compositions andmicrostructures were investigated by X-ray diffraction (XRD) andscanning electron microscopy (SEM). Then, the as-synthesizedSialon–ZrN powders were added to the SiC particles using aphenolic resin as a binder. Free-fired refractories were thenprepared and their physical properties were tested. In order tomake the as-prepared SiC–Sialon–ZrN free-fired refractories meetthe application requirements for iron-making, we further performedBF slag erosion experiments. We evaluated their BF slag erosionbehavior and investigated the mechanism underlying their BF slagresistance.

2. Experimental procedures

2.1. Synthesis of Sialon–ZrN powders from low-grade bauxiteand zirconite additives using CRN

The main raw materials employed in this study were asfollows: (i) low-grade bauxite (granularity r74 μm, chemicalcomposition (wt%): SiO2 43.55, Al2O3 40.83, Fe2O3 0.63,CaO 0.83, MgO 0.58, TiO2 1.98, K2O 0.12, Na2O 0.08 andvolatiles 11.4; from Yangquan, Shanxi Province, China),(ii) carbon coke powder (granularity r74 μm, carbon con-tent¼88 wt%; from Shanxi Xinshidai Imp. and Exp. Co., Ltd.,China), and (iii) nitrogen gas (purity 99.99%). The maincrystalline phases of the low-grade bauxite were kaolinite,halloysite, and diaspore. Zirconite was used as an additive(granularity r74 μm, its chemical composition (wt%): ZrO2

62.63, SiO2 33.17, HaO2 2.72, Al2O3 0.93, Y2O3 0.21, Fe2O3

0.12, CaO 0.09, TiO2 0.08 and Cr2O3 0.05; from ShandongProvince, China).

First, only the low-grade bauxite was used as the raw material tobe converted to Sialon via CRN. In principle, stoichiometricamounts of the ingredients could be used assuming that SiO2 inthe low-grade bauxite was reduced by carbon and totallytransformed to Si3N4, which subsequently reacted with Al2O3 toform Sialon. Our previous studies on Sialon powder synthesis,however, found that Al2O3 and X-Sialon were formed in theproduct if stoichiometric amounts of carbon coke were used. Incontrast, 15R-Sialon and SiC would form if ˃1.1 times thestoichometric amount was used. Therefore, the carbon cock contentwas fixed in this work at 1.1 times the required stoichiometricvalue (i.e., excess 10% carbon).The mixtures were milled with ethanol (A.R., Beijing

Chemical Works) in a plastic jar with agate balls for 24 h.This was followed by drying and sieving. Samples of themixed powder (1.0 g) were die-pressed at 30 MPa to formspecimens 10 mm in diameter. The specimens were placed in agraphite crucible and heated in a furnace at 1300, 1400, 1500and 1600 1C for 3 h under 0.1 MPa of nitrogen. The firedsamples were then furnace-cooled to room temperature.Based on the CRN of low-grade bauxite outlined above, we

investigated the effect of the zirconite content on the synthe-sized products by using 0%, 5%, 10%, 15% and 20% (bymass) of the additive in the CRN of the low-grade bauxite atthe optimal temperature of 1600 1C. The mixtures of low-gradebauxite, zirconite and carbon cock were mixed, milled, andpassed through a 200-mesh sieve. Then they were placed in agraphite crucible which was put in flowing nitrogen in afurnace at 1600 1C for 5 h.After the in situ nitridation reaction, the crystalline phases

were examined using an XRD (XD-3, China). Cu Kα1radiation (λ¼1.5406 Å) was employed with a step of 0.021(2θ) and scanning rate of 41 min�1. The morphology of theproducts was observed using an SEM (JSM-6460, Japan)equipped with an energy dispersive spectroscopy (EDS)detector (INCA, Oxford Instrument).

2.2. Preparation and characterization of the SiC–Sialon–ZrNfree-fired refractories

SiC particles (120 mesh, Pingluo Binhe SiC Co., Ltd,Ningxia Province, China), the as-synthesized Sialon–ZrNpowders (240 mesh), and phenolic resin (PF-5405, SinosteelLuoyang Institute of Refractories Research Co., Ltd., HenanProvince, China) were used as the basic raw materials toproduce SiC–Sialon–ZrN free-fired refractories samples. Thepreparation procedures for the SiC–Sialon–ZrN firing-freerefractories are as follows.First, the phenolic resin was prepared as a binder. Phenolic

resin solution, alcohol, and hexamine (mass ratios 72:6:22)were weighed and then the hexamine powder was added to theethanol with continuous stirring using a glass rod until noparticles remained visible. Finally, this solution was pouredinto the phenolic resin solution and stirred until a homoge-neous solution was formed.SiC particles and the as-synthesized Sialon–ZrN powders

(% mass ratio is 46:46) were weighed and mixed with phenolic

Fig. 1. XRD patterns of the CRN synthesis products from low-grade bauxite(with 10% excess carbon) at different temperatures for 3 h: (a) 1300 1C, (b)1400 1C, (c) 15001C, and (d) 1600 1C.

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resin (8%) in a granulator. The samples were denoted as S1 toS5 according to the zirconite content used to synthesizeSialon–ZrN powders (0, 5, 10, 15, and 20 wt%, respectively).The mixtures were put into a large tray and left at roomtemperature for 12 h and then dried in an oven at 65 1C for 4 hin order to make the alcohol volatile. Subsequently, themixtures were pressed into compact bars (45 mm� 6 mm� 6mm) under a pressure of 40 MPa for 30 s. Then, the greenbodies were compressed using a cold isostatic pressure of150 MPa for 60 s. The bars were dried at 60 1C for 2 h, 90 1Cfor 2 h, 120 1C for 2 h and 150 1C for 8 h in a drying oven.Then, they were placed in a corundum crucible and heated in aflow of nitrogen (purity 99.99%) in a graphite furnace at 1000or 1500 1C for 3 h. Carbon coke powder was chosen as thepacking powder in order to provide the slightly reducingatmosphere required for the final product bars. All the sampleswere allowed to cool naturally.

The product's phases were examined via XRD (XD-3,China), using Cu Kα1 radiation (λ¼1.5406 Å) with a step of0.021 (2θ) and 41 min�1 canning rate. The bulk density andapparent porosity were measured in water according to theArchimedes’ principle. The bending strength was determinedusing a conventional three-point bending method. The micro-structure and morphology of the products were observed usingSEM (JEOL JSM-6460, Japan) equipped with an EDS detector(INCA, Oxford Instruments).

2.3. BF slag erosion tests

Many different methods have been used to simulate the BF slagresistance of refractory materials. Here, the slag resistance of theSiC–Sialon–ZrN free-fired refractories was investigated using astatic crucible test. This is a popular method as it is simple toperform and because many samples can be tested in a short periodof time. Crucible samples of the SiC–Sialon–ZrN free-firedrefractories were formed at a pressure of 200 MPa. Each cruciblemeasured ⌀ 50 mm� 40 mm and had an inner hole measuring ⌀20 mm� 25 mm. Each was dried at 120 1C. The crucible samplesfilled with BF slag (6 g) were put into a high temperature furnaceand heated for 3 h at 1500 1C under a simulated slightly reducingatmosphere. The chemical compositions of the BF slag were (mass%) SiO2 36.40, CaO 38.23, Al2O3 7.96, MgO 7.63, Fe2O3 1.45, F25.55 and others 5.81. After the test, the crucible samples werecooled down to room temperature and cross-sectioned using adiamond saw. The cross-sectioned profiles were photographed andtheir erosion thicknesses were measured. Refractory samples fromthe infiltrated region and refractory-slag interface were analyzedusing XRD, SEM and EDS.

3. Results and discussion

3.1. The phase compositions and microstructures

3.1.1. Products synthesized from low-grade bauxiteThe XRD patterns from CRN samples formed with 10%

excess carbon at different temperatures are shown in Fig. 1. At1300 1C, the principal products were mullite and corundum,

suggesting that mullitization of the kaolin in the low-gradebauxite has occurred, and that the diaspore has decomposedinto corundum. In addition, diffraction peaks due to Si2N2Oappear at 1300 1C, revealing that the nitridation reaction hasstarted at this temperature. At 1400 1C, Si2N2O and corundumwere not apparent and some β-Sialon (z¼1) and X-Sialonwere present in the products. When the temperature was raisedto 1500 1C, the mullite and X phase completely disappeared,and the β-Sialon diffraction peaks were enhanced, someassociated with corundum, the content of which was decreasedupon elevating the temperature to 1600 1C. At this tempera-ture, the products were β-Sialon (z¼3), and some associatedminor phases such as 21R (SiAl6O2N6) and corundum.SEM images of the nitrided products are shown in Fig. 2. At

1300 1C, a large number of irregularly-shaped mullite particu-lates were observed in the products. At 1400 1C, a smallamount of the prism-shaped β-Sialon appeared. When thetemperature was increased to 1500 1C, some plate-like cor-undum grains were found in the products, associated withsome spherical particles. EDS, Fig. 2e, indicates that thesespherical particles were rich in Si and Fe, which, together withthe XRD result in Fig. 1c, suggesting that they belong to aFe3Si phase. At 1600 1C, a large number of columnar β-Sialongrains were formed. The above results provided a usefulfoundation upon which further studies of the CRN synthesisof Sialon–ZrN powders from low-grade bauxite and zirconitecan be based.

3.1.2. Products synthesized from low-grade bauxite andzirconiteBased on the above results, the temperature chosen to

synthesize Sialon–ZrN powders from low-grade bauxite andzirconite powders was 1600 1C. The XRD patterns of sampleswith different zirconite contents nitrided at 1600 1C are shownin Fig. 3. The products had a large number of β-Sialon peaks

Fig. 2. SEM images of the products synthesized from low-grade bauxite (10% excess carbon) at different temperatures for 3 h: (a) 1300 1C, (b) 1400 1C, (c)1500 1C, and (d) 1600 1C. EDS patterns are shown in (e) detected at the round ball marked in Fig. 2c, and (f) detected at columnar crystal marked in Fig. 2d.

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with minor peaks due to phases of 21R and corundum in thecase when the additive was absent. At 5% additive, apart fromthese phases, ZrN was found to exist in the product. Withincreasing amounts of zirconite, the ZrN diffraction peaks weresignificantly enhanced. The CRN products from low-gradebauxite and 20 wt% zirconite powders at 1600 1C werepredominantly β-Sialon and ZrN. The products’ morphologiesare indicated in Fig. 4, where granules, plate-like, and shortcolumnar grains were seen to exist, which contained Si, Al, O,N, and Zr, revealed by EDS (Fig. 4f).

3.2. Physical properties and microstructures of SiC–Sialon–ZrN free-fired refractories

3.2.1. Physical propertiesThe physical properties (porosity, bulk density, and linear

change) of the samples heated at different temperatures are listedin Table 1. Their bending strengths are as shown in Fig. 5. As canbe seen from Table 1 and Fig. 5, after being dried at 150 1C, boththe bulk density and bending strength of sample S4 and S5

(represent Sialon–ZrN powders synthesized using 15% and 20%of zirconite, respectively) significantly increased. In fact, thebending strength of sample S4 reached the maximum observedvalue here (36.05 MPa). After heating at 1000 1C, the linearchanges in the samples were small and negative, suggesting thatthe samples had shrunk slightly. At the same time, the bendingstrengths of the samples heated at 1000 1C were lower than thoseat 150 1C. This may reasonably be ascribed to the ignition loss ofthe phenol resin at 1000 1C, resulting in an increase in apparentlyporosity. However, at this temperature, the Sialon, ZrN and SiCdid not sinter together, leading to a reduction in strength. Thestrength of samples was improved by changing the heatingtemperature from 1000 1C to 1500 1C. This should be as a resultof sintering occurring at the higher temperatures, which therebyenhanced the strength of the material. It is seen that the bendingstrengths of the samples heated at 150 1C were greater, relativelythan those heated at 1000 1C and 1500 1C. This may be attributedto the close adhesion between the matrix and phenolic resinbinder. After drying at room temperature and 150 1C, thephenolic resin guaranteed that the as-synthesized free-fired

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refractories would be strong. At 1500 1C, the phenolic resin wasentirely burned out, but the pores were not totally removed.Consequently, the porosity was increased, though the bendingstrength was enhanced compared to heating at 1000 1C.Considering all of the results, it is believed that the sample S4constituted the optimum composition, i.e., the optimum SiC–Sialon–ZrN firing-free refractories consisted of 46% SiC, 8%phenolic resin and 46% Sialon–ZrN, in which the latter wasprepared from 20% zirconite additive.

3.2.2. Microstructure characteristicsThe microstructures of SiC–Sialon–ZrN free fired refrac-

tories after being dried at 150 1C are shown in Fig. 6. The

Fig. 4. SEM images of the CRN products from low-grade bauxite at 1600 1C usi(f) shows the EDS pattern detected at the site marked in Fig. 4e.

Fig. 3. XRD patterns of CRN products obtained from low-grade bauxite at1600 1C using different zirconite content: (a) 0, (b) 5%, (c) 10%, (d) 15%, and(e) 20%.

particles in the samples were wrapped in the phenol resin. Thiswas particularly significant in Fig. 6a. Further observationreveals that a strong interface between the Sialon–ZrNpowders and SiC particles was formed due to bonding bythe phenol resin. This is the reason a relatively high strength inthe samples was maintained after drying at 150 1C.As the highest strength occurred in S4 (and it had other

good physical properties as well), we focused on S4's micro-structure after heat treatment at different temperatures (Fig. 7).After drying at 150 1C, the Sialon–ZrN powders are wrappedby phenol resin, and the matrix closely bonded with the SiCparticles. After heating at 1000 1C, some pores and carbonnetwork structures resulted from the burning of the phenolresin could be seen, and most of the matrix was still wrapped.When the temperature was raised to 1500 1C, the phenol resinwas mostly burned away and lots of pores were present in theproduct. The Sialon grains exhibited a well-developed colum-nar morphology, as shown in Fig. 7f. Further observation ofthe fracture surface of sample S4 indicates that transgranularfracturing occurred in the columnar Sialon grains (Fig. 7g).This is beneficial as it improved the strength of the sample.This may be one reason for the results which indicate highporosity and high strength in samples heated at 1500 1C.

3.3. Performance evaluation and BF slag corrosion resistancemechanism of the SiC–Sialon–ZrN free-fired refractories

3.3.1. Slag corrosion resistance of the SiC–Sialon–ZrN free-fired refractoriesFig. 8 shows a digital images of crucible samples of SiC–

Sialon–ZrN Free-fired refractories after slag resistance testingat 1500 1C for 3 h. As can be seen from the cross-sections, thesizes and shapes of bottoms of crucible samples were mostlyunchanged after slag erosion—only the depth of the sample

ng different zirconite content: (a) 0, (b) 5%, (c) 10%, (d) 15%, and (e) 20%.

Fig. 5. The bending strengths of the SiC–Sialon–ZrN firing-free refractoriesformed at different temperatures.

Table 1The physical properties of specimens heated at different temperatures.

Temperature (1 C) Sample Porosity (%) Bulk density (g cm�3) Linear change (%)

150 S1 14.83 2.49 —

S2 17.87 2.42 —

S3 22.10 2.23 —

S4 9.75 2.44 —

S5 15.30 2.44 —

1000 S1 25.56 2.31 �0.34S2 21.74 2.42 �0.22S3 22.83 2.35 �0.89S4 21.96 2.35 �0.89S5 20.46 2.41 �1.74

1500 S1 31.02 2.25 �1.09S2 22.70 2.45 �0.22S3 25.75 2.23 �0.74S4 28.50 2.25 �0.81S5 27.25 2.30 �0.59

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penetrated by the slag was different. However, the zonesamong air, slag and sample phases were eroded by the slag toform trumpet-shapes in which the extent of erosion was veryserious. For quantitative evaluation, the slag erosion rate wasemployed to evaluate the slag resistance capability of the SiC–Sialon–ZrN free-fired refractories in this work. This wasrepresented by the thickness of the sample eroded by the slagper unit time as follows:

V ¼ dδ

dt¼ TCb�TCa

Tðmm=hÞ

where TCb is the thickness of the sample before erosion, TCa isthe thickness after erosion, and T is the erosion time.

The slag erosion rates for SiC–Sialon–ZrN free-fired refrac-tories heated at 1500 1C for 3 h are shown in Fig. 9. With theincreasing ZrN content, the slag erosion rate initially decreasedand then subsequently increased. The rate in S1 was 0.21 mm/h, while it was only 0.083 mm/h in S4. Therefore, an

appropriate ZrN content was beneficial in enhancing the slagresistance of the SiC–Sialon–ZrN free-fired refractors.

3.3.2. Thermodynamic analysis and experimental verificationThe BF slag was mainly composed of SiO2, CaO, MgO, and

Al2O3. The first three could react with β-Sialon. Li [12] hasanalyzed the thermodynamic processes occurring betweenCaO, FeO, SiO2, MgO and Si4Al2O2N6, with the followingresults:

Si4Al2O2N6þ9ðFeOÞ ¼ 9FeðlÞþ4SiO2þAl2O3þ3N2ðgÞ ð1Þ

ΔG0 ¼ �568797�234:92TðJÞSi4Al2O2N6þ9ðCaOÞ ¼ 9CaðlÞ þ4SiO2þAl2O3þ3N2ðgÞ ð2Þ

ΔG0 ¼ �4136313�1416:22TðJÞ

Si4Al2O2N6þ9ðMgOÞ ¼ 9MgðlÞ þ4SiO2þAl2O3þ3N2ðgÞ

ð3Þ

ΔG0 ¼ �3556403�1495:28TðJÞSi4Al2O2N6þ5ðSiO2Þ ¼ 9SiOðgÞ þAl2O3þ3N2ðgÞ ð4Þ

ΔG0 ¼ �4195353�1986:68TðJÞThermodynamic calculations indicated that only the stan-

dard free energy change (ΔG0) of Eq. (1) is negative in theactual process, suggesting that this reaction should be occur-ring. However, the other reactions should hardly occur evenconsidering the actual atmosphere. Therefore, CaO, MgO, andSiO2 in the BF slag directly react with β-Sialon with difficultyunder the thermodynamic conditions used. Questions about thecompositions of the samples and how they react with BF slagtherefore remain to be answered. In this work, the chamberwas not evacuated, so the surface of the sample wouldinevitably to be oxidized during the heating process. Thereactions are as follows [1]:

Si4Al2O2N6þ9=2O2ðgÞ ¼ 4SiO2ðgÞþAl2O3þ3N2ðgÞ ð5Þ

Fig. 6. SEM images of SiC–Sialon–ZrN firing-free refractories heated at 150 1C: (a) S1, (b) S2, (c) S3, (d) S4, and (e) S5.

J. Huang et al. / Ceramics International 40 (2014) 9763–9773 9769

SiCþ2O2ðgÞ ¼ SiO2þCO2ðgÞ ð6Þ

ZrNþO2ðgÞ ¼ ZrO2þN2ðgÞ ð7ÞIn order to simulate the micro-reducing atmosphere under

blast furnace conditions, the slag corrosion tests of the sampleswere preformed in the presence of carbon coke particles in thechamber, which would be oxidized to CO and/or CO2.Therefore, the β-Sialon and SiC might react with this COand/or CO2 as follows: [13]

Si4Al2O2N6þ9COðgÞ ¼ 4SiO2ðsÞþAl2O3ðsÞþ9CðsÞ þ3N2ðgÞð8Þ

SiCðsÞþ2COðgÞ ¼ SiO2ðsÞ þ3CðsÞ ð9Þ

ZrNþ2COðgÞ ¼ ZrO2þ2CðsÞþ1=2N2ðgÞ ð10Þ

SiCðsÞþ3CO2ðgÞ ¼ SiO2ðsÞþ4COðgÞ ð11Þ

Si4Al2O2N6þ9=2CO2ðgÞ ¼ 4SiO2ðsÞ þAl2O3ðsÞ þ9=2Cþ3N2ðgÞð12Þ

The as-synthesized SiC–Sialon–ZrN free-fired refractorieswere porous materials and the above reactions could alloccured at high temperatures involved under blast furnaceconditions to form SiO2 and Al2O3. When SiO2 made contactwith the slag, it would inevitably react with CaO, MgO, andR2O and form a low-melting glass phase. This could beconfirmed using phase diagrams of the CaO–Al2O3–SiO2,K2O–SiO2–Al2O3, and MgO–SiO2–Al2O3 systems [14].The formation of a low melting point glass phase would resultin penetration of the slag into the refractory samples anddissolution of the refractory samples into the slag. This isbelieved to be the main process by which the refractorysamples were eroded by BF slag. Even if the ZrN wasoxidized to ZrO2, it did not react with those oxides to form

a low melting point glass phase. Thus, adding zirconite into thelow-grade bauxite when synthesizing Sialon–ZrN powderscould enhance the slag erosion resistance of the SiC basedfire-free refractories.The products of the slag reactions in S1 and S4 (heated at

1500 1C) were analyzed by XRD and compared with theoriginal slag phase, as shown in Fig. 10. Melilite and mullitephases, as well as a lot of glass phase, existed in the originalslag. However, melilite and diopside were present in S1 (noZrN) and S4 (ZrN from 15% zirconite). Moreover, ZrO2 alsooccurred in S4, which reveals that the ZrN had been oxidizedand that the product formed, together with the eroded sample,had dissolved into the slag. In other words, ZrO2 did not reactwith other oxides in the slag to form a low melting point phase,confirming that ZrO2 could play a crucial role in BF slagerosion resistance.

3.3.3. Microstructure of the crucible sample after erosion byBF slag at 1500 1CSEM images and EDS results of different zones of crucible

sample S4 after being eroded by slag at 1500 1C are providedin Fig. 11. It can be clearly seen from the digital photo ofcrucible sample that there was a clear interface between sample(Zone 1) and slag (Zone 2). There was a circle of whitewhiskers at the slag–air–sample interface (Zone 3), and a denselayer of a shiny glass phase (Zone 4). The microstructures ofthese zones were each examined.The results from Zone 1 are shown in Fig. 11-1a and b. The

EDS result shows that the sample matrix contained Ca and Mgfrom the original slag, revealing that infiltration of the slag intothe sample had occurred. The digital photo of Zone 2 (i.e., slag,see Fig. 11-2) indicates the presence of many large pores. Suchpores are formed because the viscosity of the melting slagincreased from boiling temperature to cooler ones, whichresulted in large pores remaining in the outer regions of the

Fig. 7. SEM images of S4 heated at different temperatures: (a and b) 150 1C (c and d) 1000 1C, and (e–g) 1500 1C.

J. Huang et al. / Ceramics International 40 (2014) 9763–97739770

slag (Fig. 11-2a). Nevertheless, the central region of the slagformed a very dense glassy phase containing Ca, Mg, Si, Al andO (Fig. 11-2b and c).

The results from the slag–air–sample interface (Zone 3)indicate that erosion occurred in the sample above the slag linedue to oxidation. This damaged the surface structure of thesample and enhanced the slag's ability to erode the sample.The thickness of the slag-eroded interface region was 1 mm.

The microstructure of the slag–air–sample interface wasfurther observed by SEM. A large number of needle-likestructures were present at the interface among the three phases.The EDS results reveal that these were composed of SiO2 andshould be spearhead-shaped double-crystals of tridymite fromthe oxidation of Sialon/Si3N4 (Fig. 11-3a and b). There wasa layer of a dense shiny glass phase on the top surfaceof the crucible sample (Zone 4). Further SEM obser-

Fig. 8. Digital images of SiC–Sialon–ZrN crucibles after BF slag resistance testing at 1500 1C for 3 h.

Fig. 9. Comparison of slag erosion rate for SiC–Sialon–ZrN free firedrefractories heated at 1500 1C for 3 h.

Fig. 10. XRD patterns of slag in different samples heated at 1500 1C for 3 h:(a) the original slag, (b) slag in S1 (c) slag in S4.

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vation, shown in Fig. 11-4a and b, suggests that the surface ofthe sample is very dense. The EDS results indicate thatthe sample surface contains Si, Al, O, and a small amountof Na, but no N. These results reveal that the top of thecrucible sample was oxidized and Na2O coming from the slagwas volatilized and deposited on the top of crucible. It reactedwith SiO2 and Al2O3 formed by oxidation to generate aglass phase.

3.3.4. Corrosion mechanism of BF slag to SiC–Sialon–ZrNfree-fired refractories

Based on the above results, we can conclude that themechanism responsible for the corrosion of the refraxtoriesby the BF slag involved oxidation-erosion-dissolution-penetration. The process included two important aspects —

dissolution of the SiC-Sialon-ZrN refractories into the BF slagand penetration of the BF slag into the refractories. The detailsof the corrosion mechanism are as follows:

(1)

Dissolution of SiC–Sialon–ZrN refractories into the BF slagThe products of the residual slag in S1 and S4 (heated at

1500 1C) were analyzed by XRD and compared with theoriginal slag phase (Fig. 10). The results indicate that thecomponents of the slag residue had changed after corrosionhad occurred. The original slag consisted of melilite, mulliteand glass phase, whereas S1 (no ZrN) and S4 (ZrN

transformed from 15% zirconite) contained melilite anddiopside. Moreover, ZrO2 also occurred in sample S4,revealing that dissolution of the SiC–Sialon–ZrN refractoriesinto BF slag had occurred.

As dissolution erosion occurs on the surface of therefractory, the dissolution rate is determined by many factors,i.e., the composition of the refractory, the content andviscosity of the liquid phase, and the saturation concentrationof the refractory components in the liquid phase (whichincreases with increasing temperature). Diffusion plays animportant role in dissolution erosion of the refractory. Fick'sfirst law of diffusion reveals that the refractory-into-slagdiffusion rate is as follows [15]:

dn=dt ¼DSdn=dx ð13Þwhere dn/dt is the flux (mol/s), D is the diffusion constant, Sis the surface area, and dn/dx is the concentration gradient inthe x-direction.

Let δ represent the thickness of the diffusion layer.The dissolved concentration ns at the surface of the extendedlayer in the working lining of the refractory (which is closeto saturation concentration) decreases to the concentration nin the BF slag. Thus the concentration gradient in thediffusion layer is

dn=dx¼ ðns�nÞ=δ ð14Þ

Fig. 11. SEM images and EDS results from different zones of crucible sample S4 after being eroded by slag at 1500 1C: (1) bottom of crucible sample, (2) slag, (3)the slag–air–sample interface, and (4) the oxidized surface of the sample.

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Substituting this into Eq. (13) gives

dn=dt¼DSðns�nÞ=δ ð15ÞIn addition, the diffusion coefficient in the slag is inverselyproportional to viscosity, Dpη�1, so Eq. (15) can bewritten as

dn=dt¼ ðAS=δηÞðni�nÞ ð16Þwhere η is the viscosity of the BF slag and A is a constant.

Eq. (16) reveals that the dissolution rate of the refractoriesinto the BF slag increases if the BF slag viscosity decreases.The viscosity of the slag (η) depends on its internalcharacteristics and the temperature. If a solid phase is presentin the slag, it will reduce the erosion of the refractories. Inthis work, ZrO2 (oxidized from ZrN in the SiC–Sialon–ZrNfree-fired refractories) did not react with other oxides in theslag to form low melting point materials. Due to its highmelting point (2715 1C), ZrO2 should be a solid phase in theslag and as such it should reduce the erosion rate of the as-prepared refractories.

(2)

Penetration of the BF slag into the SiC–Sialon–ZrNrefractories

In general, a refractory material has a large number ofpores. The quantity and size of the pores have an importanteffect on the erosion-resistance of refractories because theydetermine the penetration rate of the BF slag into therefractories. The penetration rate of BF slag in a capillary(pore) is determined by the following equation [15]:

dL=Dt ¼ r2dð2sU cos θ=drL�gU sin υÞ=δη ð17Þ

where dL/dt is the penetration rate of the molten slag along thepore, r is pore radius, d is the density of the molten slag (g/cm3), η is viscosity of the molten slag, s is the surface tensionof the molten slag, θ is the wetting contact angle, υ is theinclination angle of the pore, L is the depth of the molten slagin the pore, and g is the acceleration due to gravity.Eq. (17) indicates that the penetration rate of BF slag into

the refractories mainly determined by a combination of thecharacteristics of the pores (e.g. pores radius) and the surfacetension and viscosity of the molten slag. However, theviscosity of fresh molten slag changes significantly after the

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reaction between the slag and refractories. Therefore, once therefractory is fixed, one way to prevent slag penetration andreduce erosion is to increase the viscosity of the slag. On theother hand, Chen [16] has suggested that improving thepenetration-resistance of refractories themselves and forminga high melting point composite between the refractories andslag can prevent slag penetration. From the XRD, SEM, andEDS results on the SiC–Sialon–ZrN free-fired refractories, wecan conclude that the slag-resistance penetration abilitybecame stronger in samples formed by adding ZrN. Thereason for this is that ZrN itself has high slag penetrationresistance. Even though the ZrN could be oxidized to formZrO2, the ZrO2 did not infiltrate the slag and Fe metal and thuseffectively served to improve slag resistance.

4. Conclusions

Sialon–ZrN powders were synthesized from low-gradebauxite and zirconite. The preparation, microstructures, physi-cal properties and slag-resistance behavior of SiC–Sialon–ZrNfree-fired refractories have been investigated. The resultsindicate that

(1)

Low-grade bauxite and zirconite powders were trans-formed to β-Sialon and ZrN by carbothermal reduction-nitridation at 1600 1C. The products had grainy, plate-like,and short columnar morphologies.

(2)

The as-synthesized β–Sialon–ZrN powders were used toform SiC-based free-fired refractories. The phenolic resinused created a strongly bonded interface between theSialon–ZrN matrix and SiC particles which enhanced thestrength after drying at 150 1C. After heat treatment at1000 1C, the linear change rates of the SiC–Sialon–ZrNfree-fired refractories were very small and negative,suggesting that the samples shrunk slightly. The strengthof samples was improved when the temperature waselevated from 1000 1C to 1500 1C. This is as a result ofthe sintering that occurred at high temperatures, whichthereby enhanced the strength of the refractories.

(3)

With increasing ZrN content, the slag erosion rate of theSiC–Sialon–ZrN free-fired refractories initially decreasedand subsequently increased (after being heated at 1500 1C).Therefore, an appropriate ZrN content was beneficial inenhancing the slag resistance of the SiC–Sialon–ZrN free-fired refractories. Melilite, diopside, and ZrO2 occurred inthe slag of the SiC–Sialon–ZrN free-fired refractories. Thepresence of ZrO2 (from oxidation of the ZrN) reveals thatthe ZrO2 did not react with other oxides in the slag to formlow melting point phases. Thus, it can play a crucial role inslag erosion resistance.

(4)

The corrosion mechanism of BF slag acting on the as-prepared SiC–Sialon–ZrN free-fired refractories involved

oxidation-erosion-dissolution-penetration. It included dis-solution of the SiC–Sialon–ZrN refractories into the BFslag and penetration of the BF slag into the SiC–Sialon–ZrN refractories.

Acknowledgments

This work was financially supported by National NaturalScience Foundation of China (Grant no. 51032007). Y. G. Liualso thanks the Program for New Century Excellent Talents inUniversity (Grant no. NCET-12-0951).

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