Effect of Temperature on the Properties and Microstructures of Carbon Refractories for Blast Furnace

The typical matrix SEM micrograph of the unburned specimen is illustrated in Figure 2. The matrix was mainly composed of graphite (dark contrast), corundum (~50 μm, light gray contrast), electric-calcined anthracite (light dark contrast), fine Si (white contrast), fine SiC (gray contrast), and a trace of Fe-containing silicate impurities (bright contrast). The impurities were from the electric-calcined anthracite and graphite raw materials.

A polished section of a specimen fired at 1200 °C is presented in Figure 3(a). Unreacted Si still remained in the matrix, although product phases such as SiC and SiO₂ began to form on the surface of Si, and a large hole was present in the Si particle. At 1400 °C (Figure 3(b)), only a little unreacted Si remained, which could not be detected by XRD, but more SiC whiskers formed around the Si particle or in the holes, which were the original sites occupied by Si particles. No Si was detected in specimens fired at 1600 °C (Figure 3(c)), but the SiC whiskers and more SiC particles with different sizes formed in the holes and matrix.

The fracture surface of samples fired at different temperatures was studied by SEM. A large number of ball-like substances could be easily identified in the matrix of the samples fired at 1100 °C. Only Si and O elements were detected by EDS in these substances; and, based on the XRD results, it may be concluded that these ball substances were cristobalite (Figure 4(a)). At 1200 °C, the amount of cristobalite tended to decrease and whiskers of nanosize in diameter and micrometer size in length appeared (Figure 4(b)). According to XRD and EDS analyses, such whiskers were β-SiC, which was the same as that reported in previous studies.[12] A few of the dumbbell-shaped whiskers were also found in the interior part of the matrix at 1200 °C along with some cristobalite balls and straight β-SiC whiskers. The ball-like cristobalite was most difficult to identify in samples above 1200 °C. Temperature played an important role in the growth and increase of β-SiC whiskers. At 1400 °C, the β-SiC whiskers were dendritic, interlocked with each other to form intertextures, and were uniformly distributed in the matrix (Figure 4(c)). Up to 1600 °C (Figure 4(d)), the β-SiC whiskers were straighter and their ratio of length to diameter increased significantly, whereas some of them curled and less uniformly distributed in the matrix. Similar microstructure evolution with temperature was also observed in Al₂O₃-ZrO₂-C composite refractories;[12] however, the onset temperature of phase evolution observed in this work was ~100 °C lower than that observed in previous investigations.

The evolution of phases in the fired samples not only depended on the temperature, but also took place in different locations of the same sample. The color of fracture surface of the burned samples started to change from 1200 °C. The gray-white color of the fringe layer, which was about 5 mm in thickness, changed to gray-green color at the interior at 1200 °C, and the same phenomenon was found at higher temperatures. The thickness of the fringe layer decreased to ~2 mm at 1300 °C and then remained constant with further increase in temperature. Figure 5 indicates the phase change from the outer to interior part of the sample fired at 1400 °C. Many balls with several micrometers in diameter were observed in the fringe layer, which were identified as cristobalite by a combination of energy-dispersive X-ray and XRD analysis (Figure 5(a)). The volume of balls reduced and the amount of β-SiC whiskers increased in the middle layer (Figure 5(b)). In the interior of the sample, the balls disappeared and only β-SiC whiskers were present (Figure 5(c)).

The thermodynamic analysis was undertaken to relate those phases formed in the fired specimens to the ones predicted at equilibrium. The amount of SiC added in the batch was small and thus was not considered in the calculation process. The calculation was made under the atmosphere of 3.5 × 10⁴ Pa CO and 6.5 × 10⁴ Pa N₂,[17] and the parameter α was defined as the weight ratio of atmosphere gas to carbon brick sample. At 1100 °C, Si is an unstable phase, and when α is close to zero, the amount of SiC and Al₂O₃ is high and decreases slightly with increasing α. With an increase in α to ~0.030, the amount of Al₆Si₂O₁₃, CO, and SiO suddenly increases and thereafter Al₆Si₂O₁₃ content continuously increases, whereas the CO (10^−2.2 g) and SiO (10⁻⁵ g) contents remain constant. When α is ~0.125, the amount of N₂ and Si₃N₄ suddenly increases and the Si₃N₄ content continues increasing, whereas the N₂ remains constant with further increase in α (Figure 6(a)). At 1200 °C, AlN disappears suddenly with increasing α; the amount of Al₆Si₂O₁₃ and Si₃N₄ increases abruptly when α is ~0.20 and 0.63, respectively; and the CO (10^−1.5 g) and SiO (10⁻⁴ g) contents are higher compared to those at 1100 °C (Figure 6(b)). At 1400 °C, Al₆Si₂O₁₃ and Si₃N₄ no longer exist and the amount of CO (10⁻² to 10^−0.8 g), SiO (10⁻⁴ to 10⁻³ g), and AlN suddenly increases initially and thereafter CO and SiO continuously increase, whereas AlN reduces with α (Figure 6(c)). At 1600 °C, Al₄C₃ occurs and disappears and Al₂O₃ decreases slightly as α increases, and the amount of CO (10^−1.3 to 1 g), SiO (10^−3.3 to 10⁻² g), and AlN is higher than it is at lower temperatures (Figure 6(d)).

Thermodynamic calculations indicate that at 1100 °C, when CO, N₂, and SiO levels are close to zero, SiC, corundum, and C are the main stable solid phases. In this case, Si initially reacts with C to form SiC:

Si₃N₄ is formed by the reaction of SiC with N₂ from the atmosphere, and the reaction between SiC and CO occurs to form SiO₂, which further reacts with Al₂O₃ to form Al₆Si₂O₁₃:

With increasing temperature, the N₂ content first increases (maximum at 1200 °C) and then decreases. At 1200 °C, AlN is formed according to

Up to 1400 °C, Al₆Si₂O₁₃ and Si₃N₄ disappear according to

Thereafter, SiO₂ formed is reduced by CO:

At 1600 °C, Al₄C₃ is formed by the reaction between Al₂O₃ and CO due to the presence of an increasing amount of CO:

The CO₂ formed in the system, in turn, is reacted with C to form CO:

However, in the fired samples, SiO₂ was the only Si-containing compound detected at 1100 °C (Figure 4(a)), indicating that kinetic factors hindered the formation SiC, Si₃N₄, and Al₆Si₂O₁₃. At 1200 °C, SiO₂ and SiC shells were observed on the surface of unreacted Si (Figure 3(a)), indicating Si was oxidized by CO. SiC whiskers were also identified on the unreacted Si and in the matrix (Figures 3(a) and 4(b)). The presence of large holes in the Si particle was believed to be the result of the direct oxidation of Si and simultaneous formation and evaporation of SiO (g):

SiC whiskers may be condensed from the vapor phase reaction between CO and SiO:AlN was not found due to the kinetic factor. With increasing temperature to 1400 °C, according to Reaction [8], SiO₂ was reduced by CO and the amount of SiC formed increased correspondingly. The difference in atmosphere partial pressure from the outer part to inner part of the samples could account for the difference in the phases formed. At 1600 °C, Al₄C₃ did not form due to the kinetic factor.

The weight change of specimens and the percentage of silicon consumed during heat treatment (initial silicon minus the unreacted silicon) are presented in Figure 7. Silicon either reacted to form silicon-containing compounds or vaporized. At firing temperatures below 1400 °C, the weight of the specimens increased, but interestingly it increased only slightly between 1200 °C and 1400 °C. At temperatures above 1400 °C, it decreased with temperature; the decrease was especially large above 1500 °C. Silicon started to react at temperatures as low as 1100 °C. Approximately 95 wt pct silicon was consumed when the temperature was 1400 °C; thereafter, the silicon consumption changed only slightly. The mass loss observed below 1000 °C was due to the release of the volatile species of the binder. The increase in mass was a direct consequence of the formation of new silicon-containing phases (SiO₂ and SiC) during heat treatment, whereas the decrease in mass at temperatures above 1400 °C was a result of the volatilization of SiO (g).

The evolutions of bulk and true densities of the specimens during heat treatment are illustrated in Figure 8. It is clear that the bulk density of specimens increased with the firing temperature and reached its maximum at 1400 °C, which was 1.65 g/cm³; it then declined down to 1.58 g/cm³ at 1600 °C. However, the true density increased slightly with temperature. The trend of bulk density change was similar to that of weight change. The bulk density decrease corresponded well to the apparent porosity increase (Figure 9), which was due to the release of the volatile species and the comparatively large bulk volume change above 1400 °C. The bulk volume change was believed to be due partly to the high anisotropic thermal expansion of the graphite. From 1000 °C to 1100 °C, the apparent porosity increased and the closed porosity reduced, whereas total porosity changed accordingly, which may be accounted for by the release of the volatile species of the binder. Above 1100 °C, as a general trend, total porosity change was similar to that of apparent porosity, i.e., first declined and then increased, reaching a minimum at 1400 °C, while the closed porosity changed only slightly (2.8 to 3.0 pct). The ceramic phases formed in the specimens contributed to the apparent and total porosity decreases, while the volatilization of SiO (g) resulted in the apparent and total porosity increase.

The micropore size distribution of specimens after firing at different temperatures is shown in Figure 11. The peak corresponding to the pore size diameter (D) ≈ 2 μm was observed for the specimen fired at 1000 °C. For the specimens fired at 1200 °C, two peaks appeared: the lower one was at D ≈ 1.8 μm and the higher one at D ≈ 0.2 μm. For the specimens fired at 1400 °C, two peaks at D ≈ 0.17 μm and D ≈ 0.25 μm were observed, and at 1600 °C, two peaks at D ≈ 0.5 μm and D ≈ 0.9 μm were detected. In order to better explain the changes in micropore size distribution, the mean diameter and <1 μm micropore volume of open pores at various temperatures were calculated and are presented in Figure 12. It could be seen that the mean diameter of the pores first reduced and then increased as the firing temperature increased from 1000 °C to 1600 °C; however, the change was not as large as at temperatures below 1200 °C. The <1 μm micropore volume and mean diameter of open pores changed on the opposite directions and reached the respective extremes at 1400 °C.

The CCS of the specimens was measured and the results are given in Figure 10. The maximum strength of 39 MPa was achieved at 1200 °C to 1400 °C. It should be noted that the trend of the strength mirrored that of the weight change and the bulk density in the range of 1000 °C to 1600 °C. The results seemed to indicate a strong relationship between the strength and the mass gain on firing.

The thermal conductivity of specimens fired was measured at temperatures of 25 °C, 300 °C, 600 °C, and 800 °C, and the results are plotted against firing temperature in Figure 13. When the firing temperature was below 1400 °C, the thermal conductivity of specimens first increased and then reduced with a maximum of 15.7 W/(m·K) at 600 °C for the specimens fired at 1200 °C. Thereafter, there was a dramatic increase in thermal conductivity of up to 19 W/(m·K) at 600 °C for specimens fired at 1500 °C. The micropore structure was closely related to the thermal conductivity of specimens. Figure 14 indicates the relationship between thermal conductivity at 600 °C and <1 μm micropore volume for specimens fired at 1100 °C to 1600 °C. This shows that the <1 μm micropore volume increased while the thermal conductivity first increased and then declined in the temperature range 1000 °C to 1400 °C, and the thermal conductivity maximum was reached when the <1 μm micropore volume of open pores was 55 pct, i.e., at 1200 °C. Above 1400 °C, the <1 μm micropore volume of open pores reduced while the thermal conductivity augmented. In the temperature range of 1200 °C to 1500 °C, the change of thermal conductivity was in the direction opposite that of the <1 μm micropore volume of open pores. The results seemed to indicate that the thermal conductivity was strongly related to the micropore volume and the phase change on firing. Increasing the micropore volumes could cause a decrease in the thermal conductivity, whereas the high thermal conductivity ceramics such as SiC could lead to an increase in the thermal conductivity. At 1000 °C to 1200 °C, the >1 μm micropore volume of open pores was 45 to 70 pct and the heat was conducted by means of convection; thus, the thermal conductivity was enhanced as the temperature increased. At 1200 °C to 1400 °C, the effect of micropore on thermal conductivity became dominant, and the thermal conductivity declined as the micropore volume increased. Over 1400 °C, both the micropore and phases formed in the specimens worked together; thermal conductivity effectively improved as the micropore decreased (Figure 12), and the high thermal conductivity SiC phase content increased (Figures 1 and 4).