Changes in Hot Modulus of Rupture with the Characteristics of Aluminosilicate Refractories

Transcription

1 Changes in Hot Modulus of Rupture with the Characteristics of Aluminosilicate Refractories K. Dehghani and C. Allaire CIREP-CRNF Ecole Polytechnique de Montreal CRIQ Campus 8475 Christophe Colomb St., Montreal, Quebec, Canada H2M 2N9 ABSTRACT The effects of refractory characteristics (composition, impurity, creep rate, porosity, density and microstructure) on the hot modulus of rupture (HMOR) of eight aluminosilicate refractory bricks were investigated. These bricks, varying from 42 to 51 wt.% in their alumina content, have potential application in carbon baking furnaces used by the aluminum industry. Hot modulus of rupture testing was carried out at 1200 C using a three-point bending method. HMOR values increased with increasing density and decreasing porosity of refractories. Also, there was an increase in HMOR amounts when creep rate of refractories was decreased. Besides, an increase in impurity level resulted in a decrease in HMOR of materials. In the light of microstructure, the HMOR value was high when there was no sign of infiltration of silica-rich phases into intergranular microcracks and vice versa. The formation of tiellite in TiO 2 -rich refractories most likely led to lowering HMOR level for these materials. INTRODUCTION Although the works, which have been done on cold modulus of rupture (MOR), are noticeable [1-5], much less data is available regarding the hot modulus of rupture (HMOR) of ceramics and, in particular, refractories [6, 7]. Hot modulus of rupture is a good criterion to evaluate the strength of refractories in service. Also, this property is an important index of brick quality and the best indication of the expected performance of these materials during their usage. In some aspects, this property resembles the creep property according to which the plasticity of refractory under simultaneous load and temperature is evaluated. In

2 case of refractories, when matrix melting occurs over a wide temperature range, study the modulus of rupture (MOR) versus temperature can be complicated. Therefore, investigation on the temperature-dependent of MOR is of significant importance. With increasing temperature, different refractories present various modulus and strength behaviours, depending on their characteristics. In addition to composition and chemical bonding, previous thermal history and particle sizing are among the other important parameters in this regard. For instance, among the different types of refractory, the alumina-silicazirconia ones exhibited the highest modulus of rupture [6], this can be attributed to zirconia which provides greater stability at high temperatures. As an another example, Carniglia and Barna [7] reported that almost all their investigated alumina-silica and basic refractories, showed an identical behaviour, which was a slight decrease in MOR with increasing temperature from room teperature up to 1500 C. The drop in MOR was about 30 kg/cm 2 when the temperature was changed in the mentioned range. However, it was interesting that for the 70%, 86% and 88% Al 2 O 3 materials, the MOR increased with increasing temperature up to about 1000 C. Besides, in cases of 88% and 99% Al 2 O 3, there was an abrupt drop in MOR when temperature exceeded about 1000 C. In case of a high alumina (very pure) refractory such as 99% Al 2 O 3, this drop was abnormal. By contrast, in case of aluminous castables (regular calcium- aluminate cemented castables), there was a decrease in strength up to about 1000 C, while the cementing reactions led to an increase in HMOR up to 1500 C [7]. Even for a given composition or a specific type of refractory, the different behaviours can be observed depending on the temperature. These differences are attributed, for example, to the changes that take place in modulus of elasticity (E) [7]. This, in turn, can be because of the wide range of bonding phases existed in a refractory, which significantly affect the strength of material in whole. It is also possible that a bonding phase did not provide the required strength during previous thermal treatments, e.g. sintering process, which is well enough to destroy the strength of whole matrix. There are other factors that still may significantly influence the MOR/HMOR by preventing or promoting the formation of microcracks in the refractories. The most important of these are porosity, thermal expansion coefficient, anisotropy of matrix, amount of each phase and its distribution, and firing temperature. The objective of the present work was therefore to study the effects of the characteristics of eight aluminosilicate refractory bricks on their hot modulus of rupture (HMOR).

3 EXPERIMENTAL PROCEDURE The refractory bricks were received from the manufacturer. Their compositions, obtained using X-ray fluorescence (XRF) method, are shown in Table I. In addition to HMOR values, the amounts of creep rate (taken from previous work [8]) as well as the impurity for each material are also presented for comparison in this Table. Table I - The X-ray fluorescence results of the aluminosilicate refractories investigated in the present work (wt%). Composi tion Material Number Al SiO Impurity Fe 2 O ZrO MgO CaO Na 2 O 0.10 <0.10 <0.10 <0.10 <0.10 <0.10 <0.10 <0.10 K 2 O TiO P 2 O Cr *[(ε 50 - ε 20 )/30] Creep Rate (%/hr) at 1280 C Hot Modulus of Rupture (HMOR) at 1200 C MPa *ε 50 and ε 20 are respectively the strain of materials after 50 hr and 20 hr interval during creep testing.

4 Samples for HMOR testing were taken from the as received materials with the dimensions of mm. They were then heated up to 1200 C and held for 1 hr at this temperature before running the test. The heating rate was about 5 C/min. The standard method of ASTM C [9] was used for HMOR testing. The rate of the application of the load on the samples was about 780 N/min. Using this load rating, the resulting rate of increase in bending stress for the samples was about 9 MPa/min. The modulus of rupture is then calculated as follow [9]: MOR = 3NL/2wd 2 (1) Where N is load at rupture, L distance between supports, w width of specimen and d thickness of specimen. Besides, micro-structural investigations, including X-ray elemental mapping, were carried out for all samples using scanning electron microscopy (SEM) technique. RESULTS AND DISCUSSION The most important factors affecting the HMOR here are discussed below: Composition of Refractories This property mostly deals with the purity degree of raw material used in the production of present refractories. Referring to the chemical composition of these materials (see Table I), sample 2 had almost a high level of K 2 O but the lowest level of TiO 2, this likely promoted the formation of K 2 O-SiO 2 eutectic rather than tiellite formation. The same might have happened to sample 8. As previously mentioned, K 2 O form an eutectic with silica having a low melting temperature of about 767 C [10] with a viscosity of about poises lower than a quartz liquid [11]. This readily explains the penetration of such a K 2 O-SiO 2 eutectic in intergranular microcracks of samples 2 and 8. For example, the infiltration of this silica rich phases, in case of sample 2, is presented in Figure 1. That can be the reason why sample 2 showed almost a very low HMOR amount (the lowest one after sample 5). By contrast, sample 8 exhibited a moderate HMOR value. That is possibly because, in comparison, sample 8 had the lowest amount of porosity with the highest value of density, which partly

5 counterbalanced the effect of probable formation of a K 2 O-SiO 2 eutectic in this material. Figure 1 - Example of SEM micrographs presenting the primary microstructural degradation because of the infiltration of Si-rich phases (bright regions shown by arrows) into intergranular microcracks (sample 2). On the other hand, the SEM micrograph of sample 6 proved that tiellite (42wt.% TiO wt.% Al 2 O 3 + 7wt.% Fe 2 O 3 ) was formed as individual grains, having almost a blocky or irregular shape, whilst distributed coarsely within the matrix. Therefore, the formation of tiellite might be dominated in samples 5 and 6. Although, these two samples had the highest content of K 2 O, their TiO 2 levels were also high. Thus, it is more likely that, as a result of tiellite formation, sample 5 presented the lowest HMOR quantity, while sample 6 had a very low HMOR amount. As for the sample 4, which had the lowest level of K 2 O, the microprobe analysis showed no sign of K 2 O, or phases containing K 2 O, within its matrix. That could be the reason why there was also no sign of penetration in grains of this sample, though the intergranular microcracks were observed, as shown in Figure 2. This led to the highest HMOR value obtained for sample 4, which also had an almost low porosity with the highest density. The same probably happened to sample 7: no infiltration was observed in its intergranular microvoids

6 most likely due to lack of formation of phases containing K 2 O. Because of this, sample 7 presented relatively a very high HMOR level. Figure 2 - Example of SEM micrographs indicating no sign of penetration of Si-rich phases (bright regions) into intergranular microcracks (sample 4). In summary, if the TiO 2 content was about 2 wt% or more, the formation of tiellite was dominated though the K 2 O was also high (e.g. samples 5 and 6). When the TiO 2 level was less than about 1 wt% and the K 2 O content about 0.4 wt% or more, then the formation of K 2 O-SiO 2 eutectic was most probably dominated (e.g. samples 2 and 8). If the K 2 O content was less than about 0.4 wt%, no signs of infiltration of K 2 O-SiO 2 eutectic into intergranular microcraks were observed (e.g. samples 4 and 7). Microstructure Stability Three steps were observed here regarding the losing of strength of microstructure and, therefore lowering the HMOR values. These steps are as follows: 1) formation of intergranular micocracks, 2) primary microstructural degradation due to infiltration of liquids into these intergranular microvoids, (e.g. see Fig. 1 taken from sample 2),

7 3) advance microstructural deterioration or rupture because of losing the strength of whole grains/refractories (e.g. see Fig. 3 that represents the microstructure of sample 1). Figure 3 - Example of SEM micrograph showing advance microstructural deterioration (i.e. no clear grain boundary) or rupture due to losing the strength of grain boundaries and whole matrix (sample 1). Referring to the above discussion, it was expected that sample 1 (with the most microstructural degradation) should have the lowest HMOR, however that was not the case here. Thus, the almost high HMOR of this sample can be attributed to relatively its low porosity/high density as discussed subsequently. It should be mentioned that penetration of silica-rich phases into refractories matrix could significantly influence the stability and mechanical properties of these materials. For example, it was found that the strength or creep resistance of refractory bricks decreased by factor of 3 when slag penetrated into the bricks during their application [12]. As it is obvious from Figure 4, the HMOR or strength of present materials decreased with increasing their creep rate or decreasing their creep resistance. Here, best-fit line equation was used to draw the line through the data points. Regarding the effect of matrix stability of the hot strength of refractories, many workers have reported [10, 13-15] that mullite has a very significant stability at high temperatures comparing to other phases. This can be attributed to the dense lattice of mullite due to its interlocked needles as well as to its high melting point. As a result, these significant properties of mullite can also yield a microstructure with much higher stability leading to a greater HMOR value.

8 50 40 HMOR (MPa) ,001 0,002 0,003 0,004 0,005 0,006 Creep rate (%/hr) x 10-3 Figure 4 - HMOR dependence of creep rate of refractories. Porosity and/or Density The results of apparent porosity and bulk density measurements are summarized in Table II. As it was expected, there was a good relation between the amount of porosity and HMOR. Figure 5 shows how porosity and density affect the HMOR of materials. According to this Figure, HMOR was increased with increasing the density of refractories. Although, the amount of HMOR was increased with decreasing the porosity, it seems that there is an abrupt increase in HMOR level when the latter had the value of about 17% or lower. This happened when density exceeded the amount of about 2.25 g/cm 3. Table II. Average apparent porosity and density of all materials. Apparent porosity (%) Apparent density (g/cm 3 ) Materia Materia Materia Materia Materia Materia Materia Materia

9 30 3 Porosity (%) Density (g/cm 3 ) Porosity (%) 2,7 2,4 2,1 1,8 1,5 Density (g/cm3) HMOR (MPa) Figure 5 - Changes in HMOR with porosity and density of refractories. The formation of pore networks is responsible for lowering the stability of matrix, which in turn results in decreasing the HMOR. The enlargement of pores and their coalescence can lead to a spongy like sample resulting in a very low strength. Porosity is related to the strength of refractories as follow [7]: σ f = σ o exp(-bp) (2) Where σ f is fracture stress, σ o is applied stress, b is constant and P is the pore volume fraction. For example, when the percent of void volume of a refractory was exceeded 70%, the crack length was found to be about several hundred millimetres or tens of centimetres [7]. This high volume of void can easily lead to a catastrophic failure or rupture of refractories. Matrix defects Although there may be various types of matrix defect, but only the effect of intergranular cracks/microcracks, formed in the present materials, is considered here. As discussed in the previous sections, the formation of these cracks plays an important role in the quantities of both cold modulus and hot modulus rupture of refractories. That is because the fracture toughness is affected strongly by the formation of intergranular cracks. The following equation expresses the relation between MOR and crack length [7]. 1,2

10 MOR= [2.25E/c] 1/2 (3) Where c is critical crack length and E is modulus of elasticity. However, the sharpness of the crack tip is also important and should be taken into account. Rigden et al. [1] have discussed that alkali-silica reaction is a major cause of crack formation. They concluded that these cracks resulted in a significant reduction in the modulus of rupture of concretes. Bonding Phases Characteristics Considering that impurities, having a low melting point, can act as bonding phases in case of alumina-silica refractories, therefore mostly their effects are discussed here. Thus, this effect was not investigated for each individual bonding phase when the effect of all these phases was taken into account as the effect of impurities. However, among impurities, it was found that TiO 2, K 2 O and Fe 2 O 3 were the most important phases that lowered the HMOR values of present refractories. As presented in Fig. 6, there is a trend of increasing the amount of HMOR with decreasing the level of both impurities and TiO 2 +Fe 2 O 3 though the slope in case of TiO 2 +Fe 2 O 3 is not significant. Once again, best-fit lines were plotted through the experimental results. The reason for considering the effect of TiO 2 +Fe 2 O 3 is that the effect of TiO 2 is pronounced in the presence of Fe 2 O 3, and vice versa, due to the enhancement of tiellite formation Impurity (wt%) 9 6 Impurity 9 6 Fe 2 O 3 +TiO 2 (wt%) Fe2O3+TiO HMOR (MPa) Figure 6 - Effects of impurity and (TiO 2 +Fe 2 O 3 ) on the HMOR amounts.

11 A large volume of lower-melting bonding phases, located between the grains of the major phases, promotes bulk melting. This in turn enhances the plastic deformation of matrix at high temperatures. The influence of, for example, a glassy phase on HMOR can be explained in this way. At room temperature, the refractories are stronger with higher strength. That is because at low temperature the glass bond is brittle and fracture would occur without significant yielding. As a result, the fracture toughness would be low. When the temperature is raised to an intermediate level, the glass softens and degrades the strength of refractories. Consequently, the glassy bond can withstand some yielding, and therefore the fracture toughness may increase somewhat. However, at high temperatures corresponding to real application of refractories, the very low viscosity of glassy phases results in rapid crack propagation and very low fracture toughness or HMOR. CONCLUSIONS 1. The most detrimental impurities identified in this work were TiO 2 and K 2 O. For the samples containing about 2 wt% TiO 2 and/or 0.4 wt% K 2 O, the HMOR was 18 MPa, here. There was an exception for sample 4 (having 2.18 wt% TiO 2 ), most likely due to its lowest K 2 O and SiO 2 contents as well as its highest density. 2. Three steps were recognized in terms of microstructural degradation during high temperature treatments of refractories: a) formation of intergranular micocracks, b) primary microstructural degradation due to infiltration of liquids into these intergranular micovoids, and c) advanced microstructural deterioration or rupture because of losing the strength of whole grains/refractories. 3. Although, the amount of HMOR was increased with decreasing the porosity, there was an abrupt increase in HMOR level when the latter had the value of about 17% or lower. ACKNOWLEDGMENTS The authors would like to thank Aluminerie de Bécancour (A.B.I) for their financial support during the realization of this work. Dr. S. Afshar is also acknowledged for his help during the course of this research.

12 References 1. S.R. Rigden, Y. Majlesi and E. Burley, Magazine of Concrete Research, vol. 47, 1995, p R. Fu and T.Y. Zhang, J. Am. Ceram. Soc., vol. 81, 1998, p H.R. Hwang and R.Y. Lee, J. of Materials Science, vol. 31, 1996, p Z.P. Bazant and Z. Li, J. of structural Engineering, vol. 121, 1995, p N.K. Bairagi and N.S. Dubal, Indian Concrete J., August, 1996, p M. J. Hendricks, M.J.P. Wang, R.A. Filbrun and D.K. Well, Modern Casting, vol. 89, 1999, p S.C. Carniglia and G.L. Barna, Handbook of Industrial Refractories Technology, Noyes Publication, New Jersey, USA, K. Dehghani and C. Allaire, "Microstructure Dependence of Creep in Aluminosilicate Refractories Used in Aluminum Industry: part 1: Impurity", submitted to Aluminum Transactions, Annual handbook of ASTM standard, C , p. 120, H. Schneider and A. Majdic, Ber Dt. Keram. Ges.,vol. 57, 1980, p B.A. Weichula and A.L. Roberts, Trans. Bri. Ceram. Soc., vol. 51, 1952, p S.M. Wiederhorn, R.F. Karuse and J. Sun, Am. Ceram. Soc. Bull., vol. 67, 1988, p H. Schneider, K. Okada and J.A. Pask, Mullite and Mullite Ceramics, John Wiley and Sons, Chichester, UK, W.W. Wright and M.S.J. Gani, J. Aust. Ceram. Soc., vol. 34, 1998, p K. Dehghani and C. Allaire, "Parameters Affecting the Creep Behavior of Aluminosilicate Refractories Used in Aluminum Industry", submitted to Aluminum Transactions, 2001.