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Influence of spherical porous aggregate content on microstructures and properties of gas-permeable mullite-corundum refractories
Ceramics International 45 (2019) 17268–17275
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Influence of spherical porous aggregate content on microstructures and properties of gas-permeable mullite-corundum refractories
T
Zhenyan Zhanga, Wen Yana,b,∗, Nan Lia,∗∗, Yuanbing Lia, Wenying Zhoua, Bingqiang Hana a
The State Key Laboratory of Refractories and Metallurgy, Wuhan University of Science and Technology, Wuhan, 430081, China National-provincial Joint Engineering Research Center of High Temperature Materials and Lining Technology, Wuhan University of Science and Technology, Wuhan, 430081, China b
A R T I C LE I N FO
A B S T R A C T
Keywords: Gas-permeable mullite-corundum refractories Spherical porous aggregate content Microstructure Gas permeability Strength
In this study, six gas-permeable mullite-corundum refractories with 45–95 wt% spherical porous mullite aggregates were prepared. The influence of the spherical porous aggregate content (SPAC) on the microstructure, gas permeability, and mechanical properties was analysed by XRD, SEM, EDS and image analysis method, etc. It's found that the SPAC affected the pore characteristics and mechanical properties through changing the packing and sintering behaviours of the aggregate and matrix in the refractories. When the SPAC was 45–65 wt %, the refractories had higher strengths as well as two gas-permeable channels in the aggregate and matrix respectively. Whereas, when the SPAC increased to 75–95 wt%, the refractories had lower strength as well as one gas-permeable channel in the matrix. The optimized product contained 65 wt% spherical porous aggregate and combined a high gas permeability of 4.9 × 10−12 m2, a high compressive strength of 18.5 MPa, a good thermal shock resistance and well-distributed gas-permeable channels.
1. Introduction Gas-permeable refractories are widely used in iron and steel industry to stir molten iron and steel and remove nonferrous inclusions from them. In order to improve the service life and the gas permeability uniformity of gas-permeable refractories, some research works have been done on the pore structure, the strength and the thermal shock resistance of gas-permeable refractories [1–4]. The pore-forming methods [5–7] of the gas-permeable ceramics mainly include the direct-foaming process, the adding pore-forming agent process, and particle packing process. In these processes, the porous ceramics prepared by the direct foaming process have high porosities and gas permeabilities, but low strengths. For instance, a porous cordierite ceramic with the high porosity of 85% was prepared by the direct-foaming method, it had a high gas permeability of 1.8 × 10−11 m2, but a low compressive strength of 2.5 MPa [8]. Therefore, this method is not suitable to prepare gas-permeable refractories for iron-making and steel-making industry. The adding pore-forming agent process is one main method to prepare the gas-permeable refractories with unidirectionally aligned pores for steel making [9–12]. The pore size and the gas permeability can be controlled by changing the diameter and amount of the pore-
∗
forming agent. However, it is difficult to form micropores in the gaspermeable refractories and improve the uniformity of gas permeability, due to the poor dispersibility of the fine pore-forming agent in raw materials. To address these problems, a particle packing method is employed to prepare the gas-permeable refractories with interconnected pores [13–17]. The smaller the particle size of the raw materials, the smaller the pore size; the narrower the particle size distribution, the more uniform the pore size distribution [18–20]. For example, four gaspermeable silicon carbide ceramics with high gas permeabilities of 1.45 × 10−12 m2 to 6.53 × 10−12 m2 have been prepared by using mono-size silicon carbide particles (mesh sizes of #120, #500, #1000, and #2000)as raw materials [21]. However, the high raw material cost caused by the narrow particle size distribution of raw materials and the low thermal shock resistance caused by the uniform structure of the refractories have limited the development of the gas-permeable refractories with interconnected pores for iron making and steel making. In order to improve the thermal shock resistance of refractories, coarse particles and fine powders were often used together to prepare gaspermeable refractories [22,23]. However, the coarse particles in the traditional gas-permeable refractories are dense aggregates and hence have no permeable passage, thus resulting in a worse permeability
Corresponding author. No. 947 Heping Road, Qingshan District, Wuhan City, Hubei Province, 430081, China. Corresponding author. No. 947 Heping Road, Qingshan District, Wuhan City, Hubei Province, 430081, China. E-mail addresses: [email protected] (W. Yan), [email protected] (N. Li).
∗∗
https://doi.org/10.1016/j.ceramint.2019.05.284 Received 25 February 2019; Received in revised form 19 May 2019; Accepted 26 May 2019 Available online 27 May 2019 0272-8842/ © 2019 Elsevier Ltd and Techna Group S.r.l. All rights reserved.
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uniformity than that of refractories with uniform structures. In order to improve the permeability uniformity of the gas-permeable refractories, it is necessary to provide more gas-permeable channels throughout refractories, which could be achieved by using porous aggregates to replace dense aggregates in conventional gas-permeable refractories [24–27]. For instance, Nihat [28] used porous coarse aggregates instead of dense aggregates in structure concrete, which improved the gas permeability and the gas permeability uniformity of the structural concrete. In our early work [29–35], it was found that when porous aggregates were used instead of dense magnesia aggregates in the periclase-spinel refractory, both the strength and the thermal shock resistance of the refractory increased remarkably due to a better interface formed between the porous aggregate and the matrix. These research results mean that using the porous aggregate to replace the dense aggregate in gas-permeable refractories can not only achieve a uniform gas permeability but also obtain a higher strength and a better thermal shock resistance. However, traditional porous aggregates are usually obtained by crushing and sieving porous ceramics [36–39]. It is difficult to control the morphology and particle size distribution of the existing porous aggregates in this way, which will further lead to the instability of pore structures and properties of the gas-permeable refractories. The purpose of this work is to use spherical porous mono-size aggregates to prepare gas-permeable refractories with gas-permeable channels both in their aggregates and matrices for iron making. The aggregates used in this work were prepared according to the patent technology in our early work [40] and had a lower cost than traditional porous aggregates. In addition, the effects of the spherical porous aggregate content (SPAC) on the microstructures, gas permeabilities and mechanical properties of gas-permeable refractories were investigated. 2. Experimental procedure 2.1. Experimental materials Spherical porous mono-size aggregates (particle diameters were 4.5 ± 0.3 mm, 1.8 ± 0.3 mm, 0.7 ± 0.2 mm, respectively), kyanite powder (< 100 μm), tabular alumina powder (< 74 μm), alumina micro-powder (< 5 μm) and white clay powder (D50 = 1.727 μm) were used as raw materials, and their chemical compositions are listed in Table 1. The spherical porous aggregates were prepared by mixing, spray granulating, baking and sintering processes, using bauxite powder, clay powder and industrial alumina powder as raw materials, carboxymethyl cellulose as a binder, and nutcracker powder as a poreforming agent [40]. The spherical porous aggregates mainly consisted of 82 wt% mullite, 17 wt% corundum, and minor rutile. Their bulk density and apparent porosity were 1.54 g/cm3 and 46.4%, respectively, and their apparent morphology is shown in Fig. 1. 2.2. Preparation of samples Six mullite-corundum gas-permeable refractories were fabricated using spherical porous aggregates, kyanite powder, tabular alumina powder, alumina micro-powder and white clay powder as raw materials, and pulp waste liquid acting as a binder. First of all, an aggregate
Fig. 1. Apparent morphology of the spherical porous aggregate with a diameter of 4.5 mm.
mixture was obtained by mixing 18 wt% aggregates with a diameter of 4.5 mm, 42 wt% aggregates with a diameter of 1.8 mm, and 40 wt% aggregates with a diameter of 0.7 mm. Secondly, a powder mixture was obtained by mixing 18 wt% kyanite powder, 64 wt% corundum powder, 6 wt% alumina micro-powder and 12 wt% white clay powder. Subsequently, six aggregate and powder mixtures were obtained by mixing the aggregate mixture and the matrix mixture by adding 10 wt% pulp waste (the pulp concentration was 45 wt %); in this process, the spherical porous aggregate contents of the six mixtures were 45 wt%, 55 wt%, 65 wt%, 75 wt%, 85 wt%, and 95 wt%, respectively. Then the six homogenous mixtures were uniaxially pressed to rectangle parallelepiped samples (140 mm × 25 mm × 25 mm) at a pressure of 10 MPa for the measurement of apparent porosity, bulk density, strength, and thermal shock resistance. For the measurement of gas permeability, cylinder samples (Φ50 mm × H50mm) were prepared. Afterwards, all samples were dried at 110 °C for 24 h. Later the rectangle parallelepiped and cylinder samples were fired at 1400 °C for 3 h in air atmosphere and subsequently cooled down to room temperature. The samples were named as A45, A55, A65, A75, A85, and A95 according to their spherical porous aggregate contents. 2.3. Investigation methods The chemical compositions of the raw materials were determined by an inductively coupled plasma-atomic emission spectrometry according to the Chinese standard GB/T 6900-2006. The median particle size was measured by a Laser Size Analyzer (Mastersizer2000, Malvern Instruments Ltd, UK). The phase compositions were analysed using an X-ray diffractometer (XRD, X'pert Pro, Philips) with Cu Kα radiation (λ = 1.54187 Å). The XRD patterns were recorded in the 2θ range of 10–90° with a scanning speed of 2° per minute, and JCPDS cards no. 01083-1881, 01-083-2080, 01-082-0511 and 01-077-0441 were used for the identification of mullite, corundum, quartz and rutile phases,
Table 1 Chemical composition of raw materials (wt %). Items
SiO2
Al2O3
Fe2O3
CaO
MgO
K2 O
Na2O
TiO2
IL
Spherical porous aggregate Kyanite powder Tabular corundum powder Alumina micro-powder White clay powder
24.90 38.73 0.02 0.15 45.10
69.66 57.05 99.42 97.69 28.16
1.18 0.93 0.02 0.19 1.47
0.30 0.02 – – 0.13
0.25 0.72 – – 0.09
0.32 0.04 – 0.15 0.13
0.01 0.23 0.36 0.35 0.01
2.85 1.84 – – 1.87
– 0.32 – – 22.12
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respectively. The crystal phase relative contents were calculated using the SemiQuant Mode of the Highscore Plus software (version 3.0) based on the XRD results. The microstructures and compositions were observed by a scanning electron microscopy (SEM, ISM-6610, JEOL, Japan) with an energy dispersive spectrometer (EDS, QUANTAX200-30, BRUKER Company, Germany). The pore size distributions and cumulative pore area distributions were statistically analysed through an image analysis method [35,41]. The bulk densities and apparent porosities of the aggregates and the fired samples were measured based on the Archimedes’ principle method using water as a medium. The flexural and compressive strengths were measured by a universal testing machine (ETM, Wance, China) with a loading rate of 0.05 MPa/s and 0.5 MPa/s, respectively. The packing densities of the spherical aggregates were measured according to the Chinese standard YB/T 187–2001. Then the gap volume percentage in the packing aggregates was evaluated using equation (1): P = 1-ρ/ρa
(1)
Where P is the gap volume percentage in the packing aggregates, ρ is the bulk density of aggregates and ρa is the packing density of aggregates. The thermal shock resistance of the cylinder samples (Φ50 mm × H50 mm) was evaluated by comparing the compressive strength before and after one thermal shock cycle. The thermal shock test was carried out by heating the samples to 1100 °C for 30 min, which was followed by air quenching for 5 min. The gas permeability was determined through the Chinese standard GB/T3000-1999 and evaluated using equation (2) [17,42]:
2p1 h q μ= 1.322 × 10−8·η· 2 · v · d Δp p1 + p2
Fig. 2. Schematic model of equipment (1) for gas permeability; (2) for bubbling test.
(2)
where μ is the gas permeability (m2), h and d are the height and the diameter of the sample (m), qv is the flow rate (m3/s), Δp is the pressure drop from entrance to exit of the sample, η is the dynamic viscosity of the fluid (Pa·s), in this study η is 1.78 × 10−5 Pa·S, p1 is the absolute pressure going into the sample (Pa), p2 is the absolute pressure coming out of the sample (Pa). The gas permeability equipment used in this study is shown in Fig. 2 (1). The sample with a height of 50 mm and a diameter of 50 mm was fixed in the center of an emulsion sleeve. The gas pressure p3 > p1 ensures no leakage of the gas. In order to find out the effective ventilation channels of gaspermeable refractories under different pressure conditions, a refitted equipment based on Fig. 2(1) was adopted, as shown in Fig. 2(2). In this refitted equipment, a transparent cylinder water container was placed on the upper surface of the sample. When nitrogen passes through the specimen, bubbles will form in the water of the container. When the first bubble appeared from the matrix position or the aggregate position on the specimen surface, the corresponding gas pressure was recorded as Pm or Pa, respectively. 3. Results and analysis 3.1. Phase composition, apparent porosity and bulk density The XRD patterns of the fired samples are shown in Fig. 3. The crystal phases of the fired samples were mainly mullite and corundum, minor quartz and rutile. As the spherical porous aggregate content (SPAC) increased, the peak values of the mullite phase increased while those of corundum phase decreased, because the porous aggregates had higher mullite content than the matrix. The apparent porosities and bulk densities as a function of the SPAC are plotted in Fig. 4. With an increment of the SPAC from 45 wt% to 65 wt%, the bulk density decreased slightly from 1.94 g/cm3 to 1.88 g/
Fig. 3. XRD patterns of the fired samples.
cm3, and the apparent porosity changed little (41.6–42.7%). Whereas, with an increment of the SPAC from 65 wt% to 95 wt%, the bulk density decreased sharply from 1.88 g/cm3 to 1.51 g/cm3, and the apparent porosity increased remarkably from 41.8% to 51.6%. 3.2. Microstructures and pore characteristics Fig. 5 and Fig. 6 show the whole microstructures and magnified matrix microstructures of the fired samples. As shown in the figures, there are two kinds of pores in the samples: one was in the aggregate, and the other was in the matrix. The spherical porous aggregate had a heterogeneous pore structure with a lot of pores in the central part but
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were small, and nearly all the gaps between aggregates were filled by the matrix, which meant that the sample A55 had the tightest microstructure. In order to further investigate the pore characteristics of the fired samples, the pore size distributions as well as the cumulative pore area distributions of the aggregate and matrices in the fired samples were analysed and shown in Fig. 7. In all fired samples, multimodal modes of the pore size distribution were observed in the aggregate and matrices. With an increment of the SPAC from 45 wt% to 95 wt%, the curves of the cumulative pore area distributions shifted obviously toward the right and the median pore size increased. The median pore size of the matrices in the samples A45-A95 were 84.9 μm, 101.4 μm, 248.8 μm, 405.3 μm, 696.8 μm, and 967.6 μm, respectively. As for the aggregates, the median pore size was 86.8 μm, very close to that of the matrix in sample A45. Fig. 4. Influence of spherical porous aggregate content on bulk densities and apparent porosities of the fired samples.
3.3. Gas permeability and mechanical properties
fewer pores near the surface, while the matrices had well-distributed pores in their general structures. It should be noted that the pores in the matrix came from two aspects. One was the pore formed from the packing of fine powders, and the other was the gap between aggregates which was not filled by fine powders. As the SPAC increased, the amount of the matrix filled in the gaps decreased, and more gaps between the porous aggregates were left. As a result, the pore area and size of the sample increased obviously with an increase of the SPAC. When the SPAC was 55 wt%, the distances between adjacent aggregates
The effects of the SPAC on the gas permeability of the fired samples are shown in Fig. 8. As the SPAC increased from 45 wt% to 55 wt%, the gas permeability of the samples A45 and A55 increased slightly from 2.2 × 10−12 m2 to 2.6 × 10−12 m2. When the SPAC was increased to 65 wt%, the gas permeability of the sample increased significantly to 4.9 × 10−12 m2. Further increasing the SPAC to 95 wt%, the gas permeability of the sample increased sharply to 3.02 × 10−11 m2. The flexural and compressive strengths of the fired samples are shown in Fig. 9. As the SPAC increased from 45 wt% to 55 wt%, the flexural strength increased slightly from 7.5 MPa to 7.9 MPa, and the
Fig. 5. Microstructures of the fired samples. 17271
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Fig. 6. Magnified microstructures of the matrix in the fired samples.
compressive strength also increased from 20.5 MPa to 21.6 MPa. When the SPAC was 65 wt%, the flexural and compressive strengths were 6.4 MPa and 18.5 MPa, respectively. With a further increase of the SPAC to 95 wt%, the flexural strength declined to 1.6 MPa, and the compressive strength declined to 1.8 MPa. The compressive strengths of the samples after a thermal shock test are shown in Fig. 10. Although the strengths of the samples A45-A65 after thermal shock test had a certain loss, their strengths after the thermal shock test were still higher than those of the samples A85 and A95.
4. Discussion Based on the obtained results above, it was found that the SPAC significantly affected the pore characteristics of the fired samples, then affecting their gas permeability, strength and thermal shock resistance. The pore characteristics were mainly determined by the initial packing behavior of the aggregate and powder mixture, which set in place a microstructure trajectory that controlled the final pore characteristics of the fired samples. In order to investigate the effect of the SPAC on the packing density of the green body of the samples, it is necessary to study the packing properties of the aggregate and the matrix respectively and find out the exact weight ratio of the aggregate to the matrix in the sample which had the densest packing structure. The gap fraction in the packing aggregates calculated by equation (1) was 33.3%. When the gaps were
filled with powders which had an equal volume to the gaps, the green body of the sample had the densest packing structure [43–45]. The weight proportion of the aggregates in the green body of the sample with the densest packing structure can be calculated according to equation (3): k = ρaVa/(ρaVa+ρmVm)*100%
(3)
Where k is the weight proportion of the aggregates in the sample, ρ is the packing density (g/cm3), V is the volume (cm3), and the subscripts m and a represent the matrix and the aggregate, respectively. Here, the parameters of ρa and ρm were 1.09 g/cm3 and 1.62 g/cm3, respectively, while the ratio of Vm to Va was 33.3%. Then, based on equation (3), the weight proportion of the aggregates in the sample with the densest packing structure was 66.9%. It was very close to that in the sample A65, and thus the sample A65 had the densest packing structure. In the sintering stage, the matrices in the samples shrank due to their much higher sintering stress than the aggregates. With a decrease of the SPAC from 95% to 65%, an increasing amount of the sintered matrix was filled in the gap between the porous aggregates, then the volume and size of the pore in the matrix decreased gradually. When the SPAC decreased from 65% to 45%, the apparent porosity of the fired samples changed little, which may be due to the similar apparent porosities of their matrix and aggregate. The strength and thermal shock resistance of the fired samples mainly depend on their microstructures. Although the green body of the
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Fig. 9. Influence of spherical porous aggregate content on the flexural and compressive strengths of the fired samples.
Fig. 7. Pore size distributions and cumulative pore area distributions of the matrices in the fired samples. Fig. 10. Influence of spherical porous aggregate content on the compressive strengths of samples after thermal shock test.
Fig. 8. Influence of spherical porous aggregate content on the gas permeability of the fired samples.
sample A65 had the lowest gap fraction, it was the fired sample A55, not A65 that had the tightest microstructure, because the volume percentage of the matrix in the sample reduced remarkably due to the sintering shrinkage, as proved in Figs. 5 and 6. The tightest microstructure means the stronger bonds between aggregate and aggregate, as well as between aggregate and matrix. This was why the sample A55 had the highest strength. In addition to the strength, the gas-permeable characteristics including the gas permeability and the permeability uniformity are also
significant properties for the application of gas-permeable refractories. In present work, the gas-permeable refractories had possible gaspermeable channels either both in the matrix and in the aggregate or only in the matrix. In order to investigate the characteristics of gaspermeable channels in the six fired samples, the gas pressure Pa or Pm (Pa and Pm represented the corresponding gas pressure when the gas began to pass through the pores in the aggregate and matrix of the samples respectively) of the samples was measured and shown in Fig. 11. It can be seen from Fig. 11, when the SPAC increased from 45% to 65%, the corresponding gas pressure Pm decreased from 35.7 kPa to 19.8 kPa, while the gas pressure Pa increased slightly from 64.2 kPa to 68.7 kPa. Compared with the Pm, the corresponding gas pressure Pa was much higher. This meant that the gas can pass through the porous aggregate in the samples A45-A65 only at a pressure higher than their Pa value (64.2–68.7 kPa). When the SPAC was 75–95%, the gas pressure Pm decreased from 3.8 kPa to 1.5 kPa, but, even when the pressure was as high as 140 kPa, the gas still could not pass through the porous aggregate. The gas pressure required for gas to pass through the pore channels depends on the pore size and other parameters, as shown in equation (4) [46]: P = 2k·δ·cosθ/R
(4)
Where P is the gas pressure acting on capillary tubes, k is the coefficient of correction, δ is the coefficient of liquids surface tension, θ is the wetting angle of materials wetted by the liquid, and R is the pore size.
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Qm d = ⎛ m⎞ Qa ⎝ da ⎠ ⎜
Fig. 11. Gas pressures corresponding to the first bubble from the aggregate and matrix of fired samples.
The smaller the pore size, the greater the pressure required for gas to pass through the pore. With the increase of the SPAC from 45 wt% to 95 wt%, the median pore size of matrix increased, resulting in the reduction of Pm, which was in accordance with equation (4). This means that the effective pore size of the ventilation channel in the matrix is directly proportional to the median pore size of the matrix. However, for the sample A45, although the median pore size of the porous aggregate (86.8 μm) was close to that of the matrices (84.9 μm), the Pa value was greater than the Pm value. It seems to contradict equation (4), actually, it was caused by the difference of the pore structures between the aggregate and matrices. As can be seen from Figs. 5 and 6, the pores and gaps were well-distributed in the matrices of the fired samples, while the pore distribution in the spherical porous aggregate was not homogeneous: with large pores in its central part but small pores near the surface. As a result, the median pore size of the porous aggregate was not an effective pore size for the ventilation channel. In present work, since the pressure P was the minimum pressure required for the gas to pass through the matrix or aggregate, we set the corresponding pore size R in equation (4) as the effective pore size of the ventilation channel. In the subsequent analysis, the median pore size of the matrix was identified as the effective pore size of the ventilation channel of the matrix, because the two parameters were proportional to each other. According to equation (4), the effective gaspermeable pore size of the ventilation channel in the aggregate (referred as Ra) can be calculated based on the measured parameters Pa and the value of k*δ*cosθ (referred as constant K). In this calculation, K can be calculated firstly using the measured parameters Pm and R (median pore size) of the matrix. The calculated Ra of the samples A45A65 were 47.2 μm, 49.9 μm, and 71.7 μm, respectively. As for samples A75-A95, the calculated Ra should be less than 11.1 μm, 10.6 μm and 10.4 μm, respectively, because Pa was higher than 140 kPa. Obviously, the Ra values of the porous aggregates in the samples A75-A95 were much smaller than the median pore size of their corresponding matrices. Based on the effective pore size of the ventilation channels of the aggregate and matrices given above, the volume flow of the gas passing through the channel can be calculated through equation (5) [47]:
Q=
1 π⋅d 2⋅v 4
(5)
where Q is the volume flow of the gas in the channel, d is the effective pore size of the ventilation channel, and v is the fluid velocity. In the present study, the ratio of the volume gas flow through the matrix to that through the aggregate in a sample can be calculated through equation (6):
⎟
(6)
where Qm and Qa are the volume flow of gas through the pore channel in the matrix and aggregate, respectively, dm and da are the effective pore size of the ventilation channel in the matrix and aggregate, respectively. The larger the value of dm/da, the larger the value of Qm/Qa, and the lower the proportion of volume flow through the pore channel in the aggregate. In samples A45-A65, the value of dm/da were 1.8, 2.0 and 3.5, then the calculated value of Qm/Qa were 3.2, 4.1, and 12.0, respectively, and thus their volume flow proportions through the pore channel in the aggregate were 23.8%, 19.6%, and 7.7%, respectively. It indicated that when the gas went into the samples A45-A65, although the gas passed through both the aggregate and the matrix, more gas flow passed through the pores in the matrices rather than aggregates. As for the samples A75-A95, the values of dm/da should be larger than 36.6, 65.7 and 93.3, and thus the calculated Qm/Qa should be higher than 1.3 × 103, 4.3 × 103 and 8.7 × 103, respectively. It meant that when the SPAC was 75–95 wt%, the gas hardly passes through the porous aggregate at the pressure lower than 140 kPa because the Qm was far more than the Qa. The above research shows that the optimized product was the samples with 65 wt% spherical porous aggregate because it combined a relatively high strength, a good thermal shock resistance, and a high gas permeability. In addition, gas can pass through the pores both in the matrix and in the aggregate at the same time. Although this gaspermeable refractory with two-channel pore structure has been preliminarily prepared, the pore structure of the spherical porous aggregate needs to be further optimized to balance the gas volume flow through the porous aggregate and the matrix. 5. Conclusion In present work, six mullite-corundum gas-permeable refractories with two gas-permeable channels were prepared using spherical porous mullite particles as aggregates. The following results were achieved: 1) The SPAC strongly affected the pore characteristics of the samples by influencing the forming and sintering processes. When the SPAC was 55 wt%, the fired sample had the tightest microstructure. 2) When the SPAC was 45–55 wt%, the samples had better mechanical properties as well as gas-permeable channels both in the aggregate and in the matrix, but lower gas permeabilities (2.2–2.6 × 10−12 m2). When the SPAC was 75–95 wt%, the gas permeability increased significantly to 1.08–3.02 × 10−11 m2, but the mechanical properties were deteriorated remarkably, and the gas passed only through the pores in the matrix rather than those in the aggregate. 3) The optimized product was a sample with 65 wt% spherical porous aggregates which had a high compressive strength of 18.5 MPa, and a high gas permeability of 4.9 × 10−12 m2 as well as two gaspermeable channels. Acknowledgments The authors would like to thank the National Key R&D Program of China (Grant No. 2017YFB0310701), the Key Project of the National Natural Science Foundation of China (Grant No. U1860205) and the Key Project of Hubei Provincial Department of Education's Scientific Research Plan (Grant No. D20181106) for financially supporting this work. References
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[1] Q.H. Wang, Y.B. Li, S.J. Li, et al., Effects of nano-alumina content on the formation
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[2]
[3] [4]
[5]
[6]
[7]
[8]
[9]
[10] [11]
[12] [13] [14] [15] [16]
[17] [18]
[19]
[20]
[21]
[22] [23]
[24]
of interconnected pores in porous purging plug materials, Ceram. Int. 43 (2017) 16722–16726. B. Long, G. Xu, A. Buhr, et al., Fracture behaviour and microstructure of refractory materials for steel ladle purging plugs in the system Al2O3-MgO-CaO, Ceram. Int. 43 (2017) 9679–9685. A. Veitsch-Radex, H. Barthel, Segment purging plugs for steel treatment in ladles, Taikabutsu 48 (1996) 535–536. L.P. Fu, A. Huang, H.Z. Gu, et al., Improvement of durability of purging plugs using MgO micropowder for refining ladles, Int. J. Appl. Ceram. Technol. 13 (2016) 1104–1111. S. Corradetti, L. Biasetto, M. Innocentini, et al., Use of polymeric fibers to increase gas permeability of lanthanum carbide based targets for nuclear physics applications, Ceram. Int. 42 (2016) 17764–17772. W.L. Huo, X.Y. Zhang, Y.G. Chen, et al., Mechanical strength of highly porous ceramic foams with thin and lamellate cell wall from particle-stabilized foams, Ceram. Int. 44 (2018) 5780–5784. J.H. Eom, Y.W. Kim, I.H. Song, Effects of the initial α-SiC content on the microstructure, mechanical properties, and permeability of macroporous silicon carbide ceramics, J. Eur. Ceram. Soc. 30 (2010) 1283–1290. Y. Ma, J.H. Li, X. Wang, et al., Highly permeable macroporous cordierite ceramics with controlled microstructure produced by particle-stabilized emulsions with a reactive thermal treatment, J. Eur. Ceram. Soc. 37 (2017) 3203–3211. T. Isobe, Y. Kameshima, A. Nakajima, et al., Gas permeability and mechanical properties of porous alumina ceramics with unidirectionally aligned pores, J. Eur. Ceram. Soc. 27 (2007) 53–59. A.G. Veitsch-Radex, H. Barthel, W. Hammerer, et al., Segment purging plugs for steel treatment in ladles, Taikabutsu 48 (1996) 535–536. W.G. Yang, G.Q. Liu, F. Yang, et al., Mathematical simulation and optimization of structure and performance of air way of purging plug in ladle, Adv. Mater. Res. 997 (2014) 565–569. X.X. Huang, W.D. Xue, J. Liu, et al., Effect of slit structure of purging plug on stress field, Adv. Mater. Res. (2011) 884–887. Y. Qin, Y. Chen, G.T. Liu, Production and application of high capacity bottom blowing purging plugs for converter, Adv. Mater. Res. (2011) 3176–3181. Z.X. Li, Y.H. Wang, J.Z. Li, et al., Performance and application of corundum based dispersive purging plug, Refract. Lime 41 (2016) 14–17 22. Tatsuya Ouchi, Isao Sasaka, Jyouki Yoshitomi, et al., Development of porous material for gas purging plugs, Taikabutsu Overseas 21 (2001) 225-225. J. Kim, H. Jang-Hoon, J.H. Lee, et al., Effect of pore structure on gas permeability constants of porous alumina, Ceram. Int. 45 (2019) 5231–5239 https://doi.org/10. 1016/j.ceramint.2018.11.219. Q.H. Wang, Y.B. Li, S.J. Li, et al., Effects of critical particle size on properties and microstructure of porous, Mater. Lett. (2017) 48–51. B. Krasnyi, V. Tarasovskii, A. Krasnyi, et al., Porous permeable ceramic for fine bubble waste water aeration systems in aeration tanks, Refract. Ind. Ceram. (2011) 322–327. Q.B. Chang, Y.L. Yang, X.Z. Zhang, et al., Effect of particle size distribution of raw powders on pore size distribution and bending strength of Al2O3 microfiltration membrane supports, J. Eur. Ceram. Soc. 34 (2014) 3819–3825. Y.J. Lia, X.F. Yang, D.H. Liu, Permeability of the porous Al2O3 ceramic with bimodal pore size distribution, Ceram. Int. 45 (2019) 5952–5957 https://doi.org/10. 1016/j.ceramint.2018.12.064. B. Wang, H. Zhang, H.T. Phuong, et al., Gas permeability and adsorbability of the glass-bonded porous silicon carbide ceramics with controlled pore size, Ceram. Int. 41 (2015) 2279–2285. Q.H. Wang, Y.B. Li, S.J. Li, et al., Effects of particle size of Al(OH)3 on the properties of porous purging materials, J. Ceram. Soc. Jpn. 125 (2017) 504–508. Q.H. Wang, Y.B. Li, S.J. Li, et al., Effects of nano-alumina content on the formation of interconnected pores in porous purging plug materials, Ceram. Int. 43 (2017) 16722–16726. M. Koohmishi, G. Shafabakhsh, Drainage potential of reservoir course of porous pavement for various particle size distributions of aggregate, Transp. Geotech. 16
(2018) 63–75. [25] V. Ducman, B. Mirtic, Water vapour permeability of lightweight concrete prepared with different types of lightweight aggregates, Constr. Build. Mater. 68 (2014) 314–319. [26] J.A. Bogas, M.G. Gomes, S. Real, Capillary absorption of structural lightweight aggregate concrete, Mater. Struct. 48 (2015) 2869–2883. [27] Y. Zaetang, A. Wongsa, V. Sata, et al., Use of lightweight aggregates in pervious concrete, Constr. Build. Mater. 48 (2013) 585–591. [28] A. Nihat, A. Mohammed Layth, A. Adem, Effects of nano-silica on the gas permeability, durability and mechanical properties of high-strength lightweight concrete, Constr. Build. Mater. (2017) 17–26. [29] W. Yan, N. Li, B.Q. Han, Effects of microsilica content on the performances of lightweight castables containing porous periclase-spinel aggregates, Ceramics 54 (2010) 315–319. [30] W. Yan, N. Li, Z. Miao, et al., Influence of properties of corundum aggregates on strength and slag resistance of refractory castables, J. Ceram. Process. Res. 13 (2012) 278–282. [31] L.P. Fu, H.Z. Gu, A. Huang, et al., Slag resistance mechanism of lightweight microporous corundum aggregate, J. Am. Ceram. Soc. 98 (2015) 1658–1663. [32] W. Yan, G.Z. Wu, S.B. Ma, et al., Energy efficient lightweight periclase-magnesium alumina spinel castables containing porous aggregates for the working lining of steel ladles, J. Eur. Ceram. Soc. 38 (2018) 4276–4282. [33] Z. Chen, W. Yan, S. Schafföner, et al., Effect of SiC powder content on lightweight corundum-magnesium aluminate spinel castables, J. Alloy. Comp. 764 (2018) 210–215. [34] G.Y. Wu, W. Yan, S. Schafföner, et al., Effect of magnesium alumina spinel content of porous aggregates on cement clinker corrosion and adherence properties of lightweight periclase-spinel refractories, Constr. Build. Mater. 185 (2018) 102–109. [35] S.B. Ma, W. Yan, S. Schafföner, et al., Influence of magnesium aluminate spinel powder content on cement clinker corrosion and adherence properties of lightweight periclase-spinel refractories, Ceram. Int. 43 (2017) 17026–17031. [36] N. Li, H.Z. Gu, H.Z. Zhao, Refractories Fundamental and Technology, Metallurgical Industry Press, Beijing, 2010 (in Chinese). [37] W. Yan, N. Li, Y.W. Li, et al., Effect of particle size on microstructure and strength of porous spinel ceramics prepared by pore-forming in situ, technique, Bull. Mater. Sci. 34 (2011) 1109–1112. [38] W. Yan, X.L. Lin, J.F. Chen, et al., Effect of TiO2 addition on microstructure and strength of porous spinel (MgAl2O4) ceramics prepared from magnesite and Al (OH)3, J. Alloy. Comp. 618 (2015) 287–291. [39] W. Yan, J.F. Chen, N. Li, et al., Preparation and characterization of porous MgOAl2O3 refractory aggregates using an in-situ, decomposition pore-forming technique, Ceram. Int. 41 (2015) 515–520. [40] N. Li, A Method for Producing Spherical Alumina-Silica Refractory Aggregates, CN Patent, No 201510121077.9, 2015. [41] X.L. Lin, Y. Wan, N. Li, Phase composition and pore evolution of porous periclasespinel ceramics prepared from magnesite and Al(OH)3, Sci. Sinter. 48 (2016) 137–155. [42] GB/T3000-1999, Dense Shaped Refractory Products-Determination of Permeability to Gases, Standards Press of China, Beijing, 2010 (in Chinese). [43] H. Lee, S. Koo, Liquid permeability of packed bed with binary mixture of particles, J. Ind. Eng. Chem. 20 (2014) 1397–1401. [44] H.A.C.K. Hettiarachchi, W.K. Mampearachchi, Effect of vibration frequency, size ratio and large particle volume fraction on packing density of binary spherical mixtures, Powder Technol. 336 (2018) 150–160. [45] G. Roquier, A Theoretical packing density model (TPDM) for ordered and disordered packings, Powder Technol. 344 (2019) 343–362. [46] J. Yu, X.J. Hu, Y. Huang, A modification of the bubble-point method to determine the pore-mouth size distribution of porous materials, Separ. Purif. Technol. 70 (2010) 314–319. [47] M.J. Zhang, Thermotechnical Fundamental of Materials, Metallurgical Industry Press, Beijing, 2008 (in Chinese).
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Influence of spherical porous aggregate content on microstructures and properties of gas-permeable mullite-corundum refractories
T
Zhenyan Zhanga, Wen Yana,b,∗, Nan Lia,∗∗, Yuanbing Lia, Wenying Zhoua, Bingqiang Hana a
The State Key Laboratory of Refractories and Metallurgy, Wuhan University of Science and Technology, Wuhan, 430081, China National-provincial Joint Engineering Research Center of High Temperature Materials and Lining Technology, Wuhan University of Science and Technology, Wuhan, 430081, China b
A R T I C LE I N FO
A B S T R A C T
Keywords: Gas-permeable mullite-corundum refractories Spherical porous aggregate content Microstructure Gas permeability Strength
In this study, six gas-permeable mullite-corundum refractories with 45–95 wt% spherical porous mullite aggregates were prepared. The influence of the spherical porous aggregate content (SPAC) on the microstructure, gas permeability, and mechanical properties was analysed by XRD, SEM, EDS and image analysis method, etc. It's found that the SPAC affected the pore characteristics and mechanical properties through changing the packing and sintering behaviours of the aggregate and matrix in the refractories. When the SPAC was 45–65 wt %, the refractories had higher strengths as well as two gas-permeable channels in the aggregate and matrix respectively. Whereas, when the SPAC increased to 75–95 wt%, the refractories had lower strength as well as one gas-permeable channel in the matrix. The optimized product contained 65 wt% spherical porous aggregate and combined a high gas permeability of 4.9 × 10−12 m2, a high compressive strength of 18.5 MPa, a good thermal shock resistance and well-distributed gas-permeable channels.
1. Introduction Gas-permeable refractories are widely used in iron and steel industry to stir molten iron and steel and remove nonferrous inclusions from them. In order to improve the service life and the gas permeability uniformity of gas-permeable refractories, some research works have been done on the pore structure, the strength and the thermal shock resistance of gas-permeable refractories [1–4]. The pore-forming methods [5–7] of the gas-permeable ceramics mainly include the direct-foaming process, the adding pore-forming agent process, and particle packing process. In these processes, the porous ceramics prepared by the direct foaming process have high porosities and gas permeabilities, but low strengths. For instance, a porous cordierite ceramic with the high porosity of 85% was prepared by the direct-foaming method, it had a high gas permeability of 1.8 × 10−11 m2, but a low compressive strength of 2.5 MPa [8]. Therefore, this method is not suitable to prepare gas-permeable refractories for iron-making and steel-making industry. The adding pore-forming agent process is one main method to prepare the gas-permeable refractories with unidirectionally aligned pores for steel making [9–12]. The pore size and the gas permeability can be controlled by changing the diameter and amount of the pore-
∗
forming agent. However, it is difficult to form micropores in the gaspermeable refractories and improve the uniformity of gas permeability, due to the poor dispersibility of the fine pore-forming agent in raw materials. To address these problems, a particle packing method is employed to prepare the gas-permeable refractories with interconnected pores [13–17]. The smaller the particle size of the raw materials, the smaller the pore size; the narrower the particle size distribution, the more uniform the pore size distribution [18–20]. For example, four gaspermeable silicon carbide ceramics with high gas permeabilities of 1.45 × 10−12 m2 to 6.53 × 10−12 m2 have been prepared by using mono-size silicon carbide particles (mesh sizes of #120, #500, #1000, and #2000)as raw materials [21]. However, the high raw material cost caused by the narrow particle size distribution of raw materials and the low thermal shock resistance caused by the uniform structure of the refractories have limited the development of the gas-permeable refractories with interconnected pores for iron making and steel making. In order to improve the thermal shock resistance of refractories, coarse particles and fine powders were often used together to prepare gaspermeable refractories [22,23]. However, the coarse particles in the traditional gas-permeable refractories are dense aggregates and hence have no permeable passage, thus resulting in a worse permeability
Corresponding author. No. 947 Heping Road, Qingshan District, Wuhan City, Hubei Province, 430081, China. Corresponding author. No. 947 Heping Road, Qingshan District, Wuhan City, Hubei Province, 430081, China. E-mail addresses: [email protected] (W. Yan), [email protected] (N. Li).
∗∗
https://doi.org/10.1016/j.ceramint.2019.05.284 Received 25 February 2019; Received in revised form 19 May 2019; Accepted 26 May 2019 Available online 27 May 2019 0272-8842/ © 2019 Elsevier Ltd and Techna Group S.r.l. All rights reserved.
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uniformity than that of refractories with uniform structures. In order to improve the permeability uniformity of the gas-permeable refractories, it is necessary to provide more gas-permeable channels throughout refractories, which could be achieved by using porous aggregates to replace dense aggregates in conventional gas-permeable refractories [24–27]. For instance, Nihat [28] used porous coarse aggregates instead of dense aggregates in structure concrete, which improved the gas permeability and the gas permeability uniformity of the structural concrete. In our early work [29–35], it was found that when porous aggregates were used instead of dense magnesia aggregates in the periclase-spinel refractory, both the strength and the thermal shock resistance of the refractory increased remarkably due to a better interface formed between the porous aggregate and the matrix. These research results mean that using the porous aggregate to replace the dense aggregate in gas-permeable refractories can not only achieve a uniform gas permeability but also obtain a higher strength and a better thermal shock resistance. However, traditional porous aggregates are usually obtained by crushing and sieving porous ceramics [36–39]. It is difficult to control the morphology and particle size distribution of the existing porous aggregates in this way, which will further lead to the instability of pore structures and properties of the gas-permeable refractories. The purpose of this work is to use spherical porous mono-size aggregates to prepare gas-permeable refractories with gas-permeable channels both in their aggregates and matrices for iron making. The aggregates used in this work were prepared according to the patent technology in our early work [40] and had a lower cost than traditional porous aggregates. In addition, the effects of the spherical porous aggregate content (SPAC) on the microstructures, gas permeabilities and mechanical properties of gas-permeable refractories were investigated. 2. Experimental procedure 2.1. Experimental materials Spherical porous mono-size aggregates (particle diameters were 4.5 ± 0.3 mm, 1.8 ± 0.3 mm, 0.7 ± 0.2 mm, respectively), kyanite powder (< 100 μm), tabular alumina powder (< 74 μm), alumina micro-powder (< 5 μm) and white clay powder (D50 = 1.727 μm) were used as raw materials, and their chemical compositions are listed in Table 1. The spherical porous aggregates were prepared by mixing, spray granulating, baking and sintering processes, using bauxite powder, clay powder and industrial alumina powder as raw materials, carboxymethyl cellulose as a binder, and nutcracker powder as a poreforming agent [40]. The spherical porous aggregates mainly consisted of 82 wt% mullite, 17 wt% corundum, and minor rutile. Their bulk density and apparent porosity were 1.54 g/cm3 and 46.4%, respectively, and their apparent morphology is shown in Fig. 1. 2.2. Preparation of samples Six mullite-corundum gas-permeable refractories were fabricated using spherical porous aggregates, kyanite powder, tabular alumina powder, alumina micro-powder and white clay powder as raw materials, and pulp waste liquid acting as a binder. First of all, an aggregate
Fig. 1. Apparent morphology of the spherical porous aggregate with a diameter of 4.5 mm.
mixture was obtained by mixing 18 wt% aggregates with a diameter of 4.5 mm, 42 wt% aggregates with a diameter of 1.8 mm, and 40 wt% aggregates with a diameter of 0.7 mm. Secondly, a powder mixture was obtained by mixing 18 wt% kyanite powder, 64 wt% corundum powder, 6 wt% alumina micro-powder and 12 wt% white clay powder. Subsequently, six aggregate and powder mixtures were obtained by mixing the aggregate mixture and the matrix mixture by adding 10 wt% pulp waste (the pulp concentration was 45 wt %); in this process, the spherical porous aggregate contents of the six mixtures were 45 wt%, 55 wt%, 65 wt%, 75 wt%, 85 wt%, and 95 wt%, respectively. Then the six homogenous mixtures were uniaxially pressed to rectangle parallelepiped samples (140 mm × 25 mm × 25 mm) at a pressure of 10 MPa for the measurement of apparent porosity, bulk density, strength, and thermal shock resistance. For the measurement of gas permeability, cylinder samples (Φ50 mm × H50mm) were prepared. Afterwards, all samples were dried at 110 °C for 24 h. Later the rectangle parallelepiped and cylinder samples were fired at 1400 °C for 3 h in air atmosphere and subsequently cooled down to room temperature. The samples were named as A45, A55, A65, A75, A85, and A95 according to their spherical porous aggregate contents. 2.3. Investigation methods The chemical compositions of the raw materials were determined by an inductively coupled plasma-atomic emission spectrometry according to the Chinese standard GB/T 6900-2006. The median particle size was measured by a Laser Size Analyzer (Mastersizer2000, Malvern Instruments Ltd, UK). The phase compositions were analysed using an X-ray diffractometer (XRD, X'pert Pro, Philips) with Cu Kα radiation (λ = 1.54187 Å). The XRD patterns were recorded in the 2θ range of 10–90° with a scanning speed of 2° per minute, and JCPDS cards no. 01083-1881, 01-083-2080, 01-082-0511 and 01-077-0441 were used for the identification of mullite, corundum, quartz and rutile phases,
Table 1 Chemical composition of raw materials (wt %). Items
SiO2
Al2O3
Fe2O3
CaO
MgO
K2 O
Na2O
TiO2
IL
Spherical porous aggregate Kyanite powder Tabular corundum powder Alumina micro-powder White clay powder
24.90 38.73 0.02 0.15 45.10
69.66 57.05 99.42 97.69 28.16
1.18 0.93 0.02 0.19 1.47
0.30 0.02 – – 0.13
0.25 0.72 – – 0.09
0.32 0.04 – 0.15 0.13
0.01 0.23 0.36 0.35 0.01
2.85 1.84 – – 1.87
– 0.32 – – 22.12
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respectively. The crystal phase relative contents were calculated using the SemiQuant Mode of the Highscore Plus software (version 3.0) based on the XRD results. The microstructures and compositions were observed by a scanning electron microscopy (SEM, ISM-6610, JEOL, Japan) with an energy dispersive spectrometer (EDS, QUANTAX200-30, BRUKER Company, Germany). The pore size distributions and cumulative pore area distributions were statistically analysed through an image analysis method [35,41]. The bulk densities and apparent porosities of the aggregates and the fired samples were measured based on the Archimedes’ principle method using water as a medium. The flexural and compressive strengths were measured by a universal testing machine (ETM, Wance, China) with a loading rate of 0.05 MPa/s and 0.5 MPa/s, respectively. The packing densities of the spherical aggregates were measured according to the Chinese standard YB/T 187–2001. Then the gap volume percentage in the packing aggregates was evaluated using equation (1): P = 1-ρ/ρa
(1)
Where P is the gap volume percentage in the packing aggregates, ρ is the bulk density of aggregates and ρa is the packing density of aggregates. The thermal shock resistance of the cylinder samples (Φ50 mm × H50 mm) was evaluated by comparing the compressive strength before and after one thermal shock cycle. The thermal shock test was carried out by heating the samples to 1100 °C for 30 min, which was followed by air quenching for 5 min. The gas permeability was determined through the Chinese standard GB/T3000-1999 and evaluated using equation (2) [17,42]:
2p1 h q μ= 1.322 × 10−8·η· 2 · v · d Δp p1 + p2
Fig. 2. Schematic model of equipment (1) for gas permeability; (2) for bubbling test.
(2)
where μ is the gas permeability (m2), h and d are the height and the diameter of the sample (m), qv is the flow rate (m3/s), Δp is the pressure drop from entrance to exit of the sample, η is the dynamic viscosity of the fluid (Pa·s), in this study η is 1.78 × 10−5 Pa·S, p1 is the absolute pressure going into the sample (Pa), p2 is the absolute pressure coming out of the sample (Pa). The gas permeability equipment used in this study is shown in Fig. 2 (1). The sample with a height of 50 mm and a diameter of 50 mm was fixed in the center of an emulsion sleeve. The gas pressure p3 > p1 ensures no leakage of the gas. In order to find out the effective ventilation channels of gaspermeable refractories under different pressure conditions, a refitted equipment based on Fig. 2(1) was adopted, as shown in Fig. 2(2). In this refitted equipment, a transparent cylinder water container was placed on the upper surface of the sample. When nitrogen passes through the specimen, bubbles will form in the water of the container. When the first bubble appeared from the matrix position or the aggregate position on the specimen surface, the corresponding gas pressure was recorded as Pm or Pa, respectively. 3. Results and analysis 3.1. Phase composition, apparent porosity and bulk density The XRD patterns of the fired samples are shown in Fig. 3. The crystal phases of the fired samples were mainly mullite and corundum, minor quartz and rutile. As the spherical porous aggregate content (SPAC) increased, the peak values of the mullite phase increased while those of corundum phase decreased, because the porous aggregates had higher mullite content than the matrix. The apparent porosities and bulk densities as a function of the SPAC are plotted in Fig. 4. With an increment of the SPAC from 45 wt% to 65 wt%, the bulk density decreased slightly from 1.94 g/cm3 to 1.88 g/
Fig. 3. XRD patterns of the fired samples.
cm3, and the apparent porosity changed little (41.6–42.7%). Whereas, with an increment of the SPAC from 65 wt% to 95 wt%, the bulk density decreased sharply from 1.88 g/cm3 to 1.51 g/cm3, and the apparent porosity increased remarkably from 41.8% to 51.6%. 3.2. Microstructures and pore characteristics Fig. 5 and Fig. 6 show the whole microstructures and magnified matrix microstructures of the fired samples. As shown in the figures, there are two kinds of pores in the samples: one was in the aggregate, and the other was in the matrix. The spherical porous aggregate had a heterogeneous pore structure with a lot of pores in the central part but
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were small, and nearly all the gaps between aggregates were filled by the matrix, which meant that the sample A55 had the tightest microstructure. In order to further investigate the pore characteristics of the fired samples, the pore size distributions as well as the cumulative pore area distributions of the aggregate and matrices in the fired samples were analysed and shown in Fig. 7. In all fired samples, multimodal modes of the pore size distribution were observed in the aggregate and matrices. With an increment of the SPAC from 45 wt% to 95 wt%, the curves of the cumulative pore area distributions shifted obviously toward the right and the median pore size increased. The median pore size of the matrices in the samples A45-A95 were 84.9 μm, 101.4 μm, 248.8 μm, 405.3 μm, 696.8 μm, and 967.6 μm, respectively. As for the aggregates, the median pore size was 86.8 μm, very close to that of the matrix in sample A45. Fig. 4. Influence of spherical porous aggregate content on bulk densities and apparent porosities of the fired samples.
3.3. Gas permeability and mechanical properties
fewer pores near the surface, while the matrices had well-distributed pores in their general structures. It should be noted that the pores in the matrix came from two aspects. One was the pore formed from the packing of fine powders, and the other was the gap between aggregates which was not filled by fine powders. As the SPAC increased, the amount of the matrix filled in the gaps decreased, and more gaps between the porous aggregates were left. As a result, the pore area and size of the sample increased obviously with an increase of the SPAC. When the SPAC was 55 wt%, the distances between adjacent aggregates
The effects of the SPAC on the gas permeability of the fired samples are shown in Fig. 8. As the SPAC increased from 45 wt% to 55 wt%, the gas permeability of the samples A45 and A55 increased slightly from 2.2 × 10−12 m2 to 2.6 × 10−12 m2. When the SPAC was increased to 65 wt%, the gas permeability of the sample increased significantly to 4.9 × 10−12 m2. Further increasing the SPAC to 95 wt%, the gas permeability of the sample increased sharply to 3.02 × 10−11 m2. The flexural and compressive strengths of the fired samples are shown in Fig. 9. As the SPAC increased from 45 wt% to 55 wt%, the flexural strength increased slightly from 7.5 MPa to 7.9 MPa, and the
Fig. 5. Microstructures of the fired samples. 17271
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Fig. 6. Magnified microstructures of the matrix in the fired samples.
compressive strength also increased from 20.5 MPa to 21.6 MPa. When the SPAC was 65 wt%, the flexural and compressive strengths were 6.4 MPa and 18.5 MPa, respectively. With a further increase of the SPAC to 95 wt%, the flexural strength declined to 1.6 MPa, and the compressive strength declined to 1.8 MPa. The compressive strengths of the samples after a thermal shock test are shown in Fig. 10. Although the strengths of the samples A45-A65 after thermal shock test had a certain loss, their strengths after the thermal shock test were still higher than those of the samples A85 and A95.
4. Discussion Based on the obtained results above, it was found that the SPAC significantly affected the pore characteristics of the fired samples, then affecting their gas permeability, strength and thermal shock resistance. The pore characteristics were mainly determined by the initial packing behavior of the aggregate and powder mixture, which set in place a microstructure trajectory that controlled the final pore characteristics of the fired samples. In order to investigate the effect of the SPAC on the packing density of the green body of the samples, it is necessary to study the packing properties of the aggregate and the matrix respectively and find out the exact weight ratio of the aggregate to the matrix in the sample which had the densest packing structure. The gap fraction in the packing aggregates calculated by equation (1) was 33.3%. When the gaps were
filled with powders which had an equal volume to the gaps, the green body of the sample had the densest packing structure [43–45]. The weight proportion of the aggregates in the green body of the sample with the densest packing structure can be calculated according to equation (3): k = ρaVa/(ρaVa+ρmVm)*100%
(3)
Where k is the weight proportion of the aggregates in the sample, ρ is the packing density (g/cm3), V is the volume (cm3), and the subscripts m and a represent the matrix and the aggregate, respectively. Here, the parameters of ρa and ρm were 1.09 g/cm3 and 1.62 g/cm3, respectively, while the ratio of Vm to Va was 33.3%. Then, based on equation (3), the weight proportion of the aggregates in the sample with the densest packing structure was 66.9%. It was very close to that in the sample A65, and thus the sample A65 had the densest packing structure. In the sintering stage, the matrices in the samples shrank due to their much higher sintering stress than the aggregates. With a decrease of the SPAC from 95% to 65%, an increasing amount of the sintered matrix was filled in the gap between the porous aggregates, then the volume and size of the pore in the matrix decreased gradually. When the SPAC decreased from 65% to 45%, the apparent porosity of the fired samples changed little, which may be due to the similar apparent porosities of their matrix and aggregate. The strength and thermal shock resistance of the fired samples mainly depend on their microstructures. Although the green body of the
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Fig. 9. Influence of spherical porous aggregate content on the flexural and compressive strengths of the fired samples.
Fig. 7. Pore size distributions and cumulative pore area distributions of the matrices in the fired samples. Fig. 10. Influence of spherical porous aggregate content on the compressive strengths of samples after thermal shock test.
Fig. 8. Influence of spherical porous aggregate content on the gas permeability of the fired samples.
sample A65 had the lowest gap fraction, it was the fired sample A55, not A65 that had the tightest microstructure, because the volume percentage of the matrix in the sample reduced remarkably due to the sintering shrinkage, as proved in Figs. 5 and 6. The tightest microstructure means the stronger bonds between aggregate and aggregate, as well as between aggregate and matrix. This was why the sample A55 had the highest strength. In addition to the strength, the gas-permeable characteristics including the gas permeability and the permeability uniformity are also
significant properties for the application of gas-permeable refractories. In present work, the gas-permeable refractories had possible gaspermeable channels either both in the matrix and in the aggregate or only in the matrix. In order to investigate the characteristics of gaspermeable channels in the six fired samples, the gas pressure Pa or Pm (Pa and Pm represented the corresponding gas pressure when the gas began to pass through the pores in the aggregate and matrix of the samples respectively) of the samples was measured and shown in Fig. 11. It can be seen from Fig. 11, when the SPAC increased from 45% to 65%, the corresponding gas pressure Pm decreased from 35.7 kPa to 19.8 kPa, while the gas pressure Pa increased slightly from 64.2 kPa to 68.7 kPa. Compared with the Pm, the corresponding gas pressure Pa was much higher. This meant that the gas can pass through the porous aggregate in the samples A45-A65 only at a pressure higher than their Pa value (64.2–68.7 kPa). When the SPAC was 75–95%, the gas pressure Pm decreased from 3.8 kPa to 1.5 kPa, but, even when the pressure was as high as 140 kPa, the gas still could not pass through the porous aggregate. The gas pressure required for gas to pass through the pore channels depends on the pore size and other parameters, as shown in equation (4) [46]: P = 2k·δ·cosθ/R
(4)
Where P is the gas pressure acting on capillary tubes, k is the coefficient of correction, δ is the coefficient of liquids surface tension, θ is the wetting angle of materials wetted by the liquid, and R is the pore size.
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Qm d = ⎛ m⎞ Qa ⎝ da ⎠ ⎜
Fig. 11. Gas pressures corresponding to the first bubble from the aggregate and matrix of fired samples.
The smaller the pore size, the greater the pressure required for gas to pass through the pore. With the increase of the SPAC from 45 wt% to 95 wt%, the median pore size of matrix increased, resulting in the reduction of Pm, which was in accordance with equation (4). This means that the effective pore size of the ventilation channel in the matrix is directly proportional to the median pore size of the matrix. However, for the sample A45, although the median pore size of the porous aggregate (86.8 μm) was close to that of the matrices (84.9 μm), the Pa value was greater than the Pm value. It seems to contradict equation (4), actually, it was caused by the difference of the pore structures between the aggregate and matrices. As can be seen from Figs. 5 and 6, the pores and gaps were well-distributed in the matrices of the fired samples, while the pore distribution in the spherical porous aggregate was not homogeneous: with large pores in its central part but small pores near the surface. As a result, the median pore size of the porous aggregate was not an effective pore size for the ventilation channel. In present work, since the pressure P was the minimum pressure required for the gas to pass through the matrix or aggregate, we set the corresponding pore size R in equation (4) as the effective pore size of the ventilation channel. In the subsequent analysis, the median pore size of the matrix was identified as the effective pore size of the ventilation channel of the matrix, because the two parameters were proportional to each other. According to equation (4), the effective gaspermeable pore size of the ventilation channel in the aggregate (referred as Ra) can be calculated based on the measured parameters Pa and the value of k*δ*cosθ (referred as constant K). In this calculation, K can be calculated firstly using the measured parameters Pm and R (median pore size) of the matrix. The calculated Ra of the samples A45A65 were 47.2 μm, 49.9 μm, and 71.7 μm, respectively. As for samples A75-A95, the calculated Ra should be less than 11.1 μm, 10.6 μm and 10.4 μm, respectively, because Pa was higher than 140 kPa. Obviously, the Ra values of the porous aggregates in the samples A75-A95 were much smaller than the median pore size of their corresponding matrices. Based on the effective pore size of the ventilation channels of the aggregate and matrices given above, the volume flow of the gas passing through the channel can be calculated through equation (5) [47]:
Q=
1 π⋅d 2⋅v 4
(5)
where Q is the volume flow of the gas in the channel, d is the effective pore size of the ventilation channel, and v is the fluid velocity. In the present study, the ratio of the volume gas flow through the matrix to that through the aggregate in a sample can be calculated through equation (6):
⎟
(6)
where Qm and Qa are the volume flow of gas through the pore channel in the matrix and aggregate, respectively, dm and da are the effective pore size of the ventilation channel in the matrix and aggregate, respectively. The larger the value of dm/da, the larger the value of Qm/Qa, and the lower the proportion of volume flow through the pore channel in the aggregate. In samples A45-A65, the value of dm/da were 1.8, 2.0 and 3.5, then the calculated value of Qm/Qa were 3.2, 4.1, and 12.0, respectively, and thus their volume flow proportions through the pore channel in the aggregate were 23.8%, 19.6%, and 7.7%, respectively. It indicated that when the gas went into the samples A45-A65, although the gas passed through both the aggregate and the matrix, more gas flow passed through the pores in the matrices rather than aggregates. As for the samples A75-A95, the values of dm/da should be larger than 36.6, 65.7 and 93.3, and thus the calculated Qm/Qa should be higher than 1.3 × 103, 4.3 × 103 and 8.7 × 103, respectively. It meant that when the SPAC was 75–95 wt%, the gas hardly passes through the porous aggregate at the pressure lower than 140 kPa because the Qm was far more than the Qa. The above research shows that the optimized product was the samples with 65 wt% spherical porous aggregate because it combined a relatively high strength, a good thermal shock resistance, and a high gas permeability. In addition, gas can pass through the pores both in the matrix and in the aggregate at the same time. Although this gaspermeable refractory with two-channel pore structure has been preliminarily prepared, the pore structure of the spherical porous aggregate needs to be further optimized to balance the gas volume flow through the porous aggregate and the matrix. 5. Conclusion In present work, six mullite-corundum gas-permeable refractories with two gas-permeable channels were prepared using spherical porous mullite particles as aggregates. The following results were achieved: 1) The SPAC strongly affected the pore characteristics of the samples by influencing the forming and sintering processes. When the SPAC was 55 wt%, the fired sample had the tightest microstructure. 2) When the SPAC was 45–55 wt%, the samples had better mechanical properties as well as gas-permeable channels both in the aggregate and in the matrix, but lower gas permeabilities (2.2–2.6 × 10−12 m2). When the SPAC was 75–95 wt%, the gas permeability increased significantly to 1.08–3.02 × 10−11 m2, but the mechanical properties were deteriorated remarkably, and the gas passed only through the pores in the matrix rather than those in the aggregate. 3) The optimized product was a sample with 65 wt% spherical porous aggregates which had a high compressive strength of 18.5 MPa, and a high gas permeability of 4.9 × 10−12 m2 as well as two gaspermeable channels. Acknowledgments The authors would like to thank the National Key R&D Program of China (Grant No. 2017YFB0310701), the Key Project of the National Natural Science Foundation of China (Grant No. U1860205) and the Key Project of Hubei Provincial Department of Education's Scientific Research Plan (Grant No. D20181106) for financially supporting this work. References
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[1] Q.H. Wang, Y.B. Li, S.J. Li, et al., Effects of nano-alumina content on the formation
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Z. Zhang, et al.
[2]
[3] [4]
[5]
[6]
[7]
[8]
[9]
[10] [11]
[12] [13] [14] [15] [16]
[17] [18]
[19]
[20]
[21]
[22] [23]
[24]
of interconnected pores in porous purging plug materials, Ceram. Int. 43 (2017) 16722–16726. B. Long, G. Xu, A. Buhr, et al., Fracture behaviour and microstructure of refractory materials for steel ladle purging plugs in the system Al2O3-MgO-CaO, Ceram. Int. 43 (2017) 9679–9685. A. Veitsch-Radex, H. Barthel, Segment purging plugs for steel treatment in ladles, Taikabutsu 48 (1996) 535–536. L.P. Fu, A. Huang, H.Z. Gu, et al., Improvement of durability of purging plugs using MgO micropowder for refining ladles, Int. J. Appl. Ceram. Technol. 13 (2016) 1104–1111. S. Corradetti, L. Biasetto, M. Innocentini, et al., Use of polymeric fibers to increase gas permeability of lanthanum carbide based targets for nuclear physics applications, Ceram. Int. 42 (2016) 17764–17772. W.L. Huo, X.Y. Zhang, Y.G. Chen, et al., Mechanical strength of highly porous ceramic foams with thin and lamellate cell wall from particle-stabilized foams, Ceram. Int. 44 (2018) 5780–5784. J.H. Eom, Y.W. Kim, I.H. Song, Effects of the initial α-SiC content on the microstructure, mechanical properties, and permeability of macroporous silicon carbide ceramics, J. Eur. Ceram. Soc. 30 (2010) 1283–1290. Y. Ma, J.H. Li, X. Wang, et al., Highly permeable macroporous cordierite ceramics with controlled microstructure produced by particle-stabilized emulsions with a reactive thermal treatment, J. Eur. Ceram. Soc. 37 (2017) 3203–3211. T. Isobe, Y. Kameshima, A. Nakajima, et al., Gas permeability and mechanical properties of porous alumina ceramics with unidirectionally aligned pores, J. Eur. Ceram. Soc. 27 (2007) 53–59. A.G. Veitsch-Radex, H. Barthel, W. Hammerer, et al., Segment purging plugs for steel treatment in ladles, Taikabutsu 48 (1996) 535–536. W.G. Yang, G.Q. Liu, F. Yang, et al., Mathematical simulation and optimization of structure and performance of air way of purging plug in ladle, Adv. Mater. Res. 997 (2014) 565–569. X.X. Huang, W.D. Xue, J. Liu, et al., Effect of slit structure of purging plug on stress field, Adv. Mater. Res. (2011) 884–887. Y. Qin, Y. Chen, G.T. Liu, Production and application of high capacity bottom blowing purging plugs for converter, Adv. Mater. Res. (2011) 3176–3181. Z.X. Li, Y.H. Wang, J.Z. Li, et al., Performance and application of corundum based dispersive purging plug, Refract. Lime 41 (2016) 14–17 22. Tatsuya Ouchi, Isao Sasaka, Jyouki Yoshitomi, et al., Development of porous material for gas purging plugs, Taikabutsu Overseas 21 (2001) 225-225. J. Kim, H. Jang-Hoon, J.H. Lee, et al., Effect of pore structure on gas permeability constants of porous alumina, Ceram. Int. 45 (2019) 5231–5239 https://doi.org/10. 1016/j.ceramint.2018.11.219. Q.H. Wang, Y.B. Li, S.J. Li, et al., Effects of critical particle size on properties and microstructure of porous, Mater. Lett. (2017) 48–51. B. Krasnyi, V. Tarasovskii, A. Krasnyi, et al., Porous permeable ceramic for fine bubble waste water aeration systems in aeration tanks, Refract. Ind. Ceram. (2011) 322–327. Q.B. Chang, Y.L. Yang, X.Z. Zhang, et al., Effect of particle size distribution of raw powders on pore size distribution and bending strength of Al2O3 microfiltration membrane supports, J. Eur. Ceram. Soc. 34 (2014) 3819–3825. Y.J. Lia, X.F. Yang, D.H. Liu, Permeability of the porous Al2O3 ceramic with bimodal pore size distribution, Ceram. Int. 45 (2019) 5952–5957 https://doi.org/10. 1016/j.ceramint.2018.12.064. B. Wang, H. Zhang, H.T. Phuong, et al., Gas permeability and adsorbability of the glass-bonded porous silicon carbide ceramics with controlled pore size, Ceram. Int. 41 (2015) 2279–2285. Q.H. Wang, Y.B. Li, S.J. Li, et al., Effects of particle size of Al(OH)3 on the properties of porous purging materials, J. Ceram. Soc. Jpn. 125 (2017) 504–508. Q.H. Wang, Y.B. Li, S.J. Li, et al., Effects of nano-alumina content on the formation of interconnected pores in porous purging plug materials, Ceram. Int. 43 (2017) 16722–16726. M. Koohmishi, G. Shafabakhsh, Drainage potential of reservoir course of porous pavement for various particle size distributions of aggregate, Transp. Geotech. 16
(2018) 63–75. [25] V. Ducman, B. Mirtic, Water vapour permeability of lightweight concrete prepared with different types of lightweight aggregates, Constr. Build. Mater. 68 (2014) 314–319. [26] J.A. Bogas, M.G. Gomes, S. Real, Capillary absorption of structural lightweight aggregate concrete, Mater. Struct. 48 (2015) 2869–2883. [27] Y. Zaetang, A. Wongsa, V. Sata, et al., Use of lightweight aggregates in pervious concrete, Constr. Build. Mater. 48 (2013) 585–591. [28] A. Nihat, A. Mohammed Layth, A. Adem, Effects of nano-silica on the gas permeability, durability and mechanical properties of high-strength lightweight concrete, Constr. Build. Mater. (2017) 17–26. [29] W. Yan, N. Li, B.Q. Han, Effects of microsilica content on the performances of lightweight castables containing porous periclase-spinel aggregates, Ceramics 54 (2010) 315–319. [30] W. Yan, N. Li, Z. Miao, et al., Influence of properties of corundum aggregates on strength and slag resistance of refractory castables, J. Ceram. Process. Res. 13 (2012) 278–282. [31] L.P. Fu, H.Z. Gu, A. Huang, et al., Slag resistance mechanism of lightweight microporous corundum aggregate, J. Am. Ceram. Soc. 98 (2015) 1658–1663. [32] W. Yan, G.Z. Wu, S.B. Ma, et al., Energy efficient lightweight periclase-magnesium alumina spinel castables containing porous aggregates for the working lining of steel ladles, J. Eur. Ceram. Soc. 38 (2018) 4276–4282. [33] Z. Chen, W. Yan, S. Schafföner, et al., Effect of SiC powder content on lightweight corundum-magnesium aluminate spinel castables, J. Alloy. Comp. 764 (2018) 210–215. [34] G.Y. Wu, W. Yan, S. Schafföner, et al., Effect of magnesium alumina spinel content of porous aggregates on cement clinker corrosion and adherence properties of lightweight periclase-spinel refractories, Constr. Build. Mater. 185 (2018) 102–109. [35] S.B. Ma, W. Yan, S. Schafföner, et al., Influence of magnesium aluminate spinel powder content on cement clinker corrosion and adherence properties of lightweight periclase-spinel refractories, Ceram. Int. 43 (2017) 17026–17031. [36] N. Li, H.Z. Gu, H.Z. Zhao, Refractories Fundamental and Technology, Metallurgical Industry Press, Beijing, 2010 (in Chinese). [37] W. Yan, N. Li, Y.W. Li, et al., Effect of particle size on microstructure and strength of porous spinel ceramics prepared by pore-forming in situ, technique, Bull. Mater. Sci. 34 (2011) 1109–1112. [38] W. Yan, X.L. Lin, J.F. Chen, et al., Effect of TiO2 addition on microstructure and strength of porous spinel (MgAl2O4) ceramics prepared from magnesite and Al (OH)3, J. Alloy. Comp. 618 (2015) 287–291. [39] W. Yan, J.F. Chen, N. Li, et al., Preparation and characterization of porous MgOAl2O3 refractory aggregates using an in-situ, decomposition pore-forming technique, Ceram. Int. 41 (2015) 515–520. [40] N. Li, A Method for Producing Spherical Alumina-Silica Refractory Aggregates, CN Patent, No 201510121077.9, 2015. [41] X.L. Lin, Y. Wan, N. Li, Phase composition and pore evolution of porous periclasespinel ceramics prepared from magnesite and Al(OH)3, Sci. Sinter. 48 (2016) 137–155. [42] GB/T3000-1999, Dense Shaped Refractory Products-Determination of Permeability to Gases, Standards Press of China, Beijing, 2010 (in Chinese). [43] H. Lee, S. Koo, Liquid permeability of packed bed with binary mixture of particles, J. Ind. Eng. Chem. 20 (2014) 1397–1401. [44] H.A.C.K. Hettiarachchi, W.K. Mampearachchi, Effect of vibration frequency, size ratio and large particle volume fraction on packing density of binary spherical mixtures, Powder Technol. 336 (2018) 150–160. [45] G. Roquier, A Theoretical packing density model (TPDM) for ordered and disordered packings, Powder Technol. 344 (2019) 343–362. [46] J. Yu, X.J. Hu, Y. Huang, A modification of the bubble-point method to determine the pore-mouth size distribution of porous materials, Separ. Purif. Technol. 70 (2010) 314–319. [47] M.J. Zhang, Thermotechnical Fundamental of Materials, Metallurgical Industry Press, Beijing, 2008 (in Chinese).
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