In-Situ Mullite Zirconia Composites from Kaolin

Transcription

1 (a) (b) (c) (d) In-Situ Mullite Zirconia Composites from Kaolin Greater content in a mullite composite made from kaolin produced increased density and higher flexural strength. (e) S.H. Kenawy, M. Awaad and H. Awad Mullite ceramics continue to be significant in the development of traditional and advanced ceramics. Mullite is the only stable crystalline phase in the aluminosilicate system of chemical composition that ranges from 3 2SiO 2 to approximately 2 SiO 2. 1 The incongruent melting behavior of mullite now is widely accepted. Because of its peculiar physical properties low thermal expansion ( / C), creep resistance, low thermal conductivity and high chemical stability mullite is considered as a compound with high technological interest. 1 However, low mechanical strength at room temperature may limit its use. Also, mullite is difficult to consolidate to fully dense, single-phase bodies because of the high activation energy for ion diffusion through the mullite lattice. One of the basic approaches to overcome this problem is the incorporation of a second phase, e.g.,. 2 Mullite composites are widely used as components in ceramic materials because of their excellent resistance to creep, spalling and high mechanical strength. 3 Mullite multicomposites are widely used in forehearth feeders and glassmelting furnaces, especially in glass contact zones, such as plungers, spouts or tubes. The widespread use of this (Top) Fig. 4 SEMs of the composite samples fired at 1600 C for 2 h: (a) MZ0 (white areas are glassy phase, (b) MZ5, (c) MZ10, (d) MZ15 and (e) MZ20 (white areas are and dark areas are mullite) American Ceramic Society Bulletin, Vol. 85, No. 11

2 class of materials is related to its high corrosion resistance, which is a result of its microstructure. Moreover, is slightly wetted by siliceous and metallic melts. Because has a low tendency to dissolve in SiO 2, the chemical attack of the composite is comparatively low. 4 Mullite ceramics can be obtained through various processing techniques, 5 including sol gel mixing of SiO 2, and ; reaction sintering of zircon and ; or mechanical mixing of mullite and. The sol gel technique is expensive. Reaction sintering of zircon and may lead to excessive undesired TG Temperature ( C) Fig. 1 Differential thermal analysis and thermogravimetric analysis of starting materials. glassy phases because of the presence of additives that are arbitrarily added to promote the reaction sintering processes by enabling reaction and densification simultaneously. 6 A study was conducted to prepare and characterize in-situ mullite reinforced with composites using Egyptian kaolin and micron-sized. Addition of to the mixture was conducted to reach the stoichiometric composition of mullite. The object was to produce the least-expensive high-temperature ceramic bodies for use in a variety of refractory and ceramic applications. Materials The major raw materials used in the study were kaolin from the Tieh area of South Sinai, ultrapure and calcined (Alcoa, Pittsburgh, Pa.), ultrapure and fine 3-mol%-Y 2 -stabilized tetragonal- (t-) polycrystals (TZP; Tosoh, Tokyo, Japan) (Table 1). Kaolin was ground and sieved to pass through a 63 µm standard sieve. A calcination process was conducted at 1000 C for 2 h to prevent the delatometric effect at 950 C for mullite. Batch compositions (Table 2) were selected in such a way that they contained 5 20 wt% and wt% (an /clay ratio of 1:1). Sample Preparation The batch design of composites contained nominal contents of 0, 5, 10, 15 and 20 wt% (MZ0, MZ5, MZ10, MZ15 and MZ20) (Table 2). The batches were attrition milled in isopropyl alcohol using 1 2 mm diameter zirconia balls as grinding media for an optimum time period of 4 h. The resulting slurries were dried at 110 C and sieved to break up agglomerates. Powdered samples were molded as cm compacts and were fabricated using uniaxial pressing at 250 MPa using a stainlesssteel die. The green compacts were fired at 1600 C with 2 h soaking time in an electric furnace using a heating rate of 300 C/h. Fired-Body Characterization The sintered samples were characterized to determine their bulk density, apparent porosity, flexural strength at room temperature, phase assembly and microstructure. The density of the sintered compacts was measured using the conventional liquid-displacement method and the Archimedes principle in water according to Standard Test Methods for Apparent Porosity, Water Absorption, Apparent Specific Gravity and Bulk Density of Burned Refractory Brick and Shapes by Boiling Water, (ASTM Designation C20. ASTM Book of Standards, ASTM International, West Conshohocken, Pa.). X-ray diffractometry (XRD) patterns of the fired samples were obtained (Model D-500, Siemens, Karlsruhe, Germany) using CuKα radiation and a graphite monochromator with a goniometric range of θ. Surface-area determination (BET) was conducted (Model ASAP200, Micromeritics Instruments GmbH, Munich, Germany) using nitrogen as absorbent gas. Particle-size distribution also was determined (Model Mastersizer 2000, Malvern Instruments, Malvern, U.K.). DTA (µv) American Ceramic Society Bulletin, Vol. 85, No

3 Table 1 Starting Material Compositions Constituent Kaolin (%) (%) (%) < SiO TiO Nil Y Nil Fe Nil 0.04 MnO Nil Nil Nil CaO Nil Nil 0.03 Na 2 O K 2 O 0.10 Nil Nil P 2 O Nil Nil Cl 0.06 Nil Nil SO Nil Nil LOI Nil Differential thermal analysis (DTA) and thermogravimetric analysis (TGA) studies of the kaolin were conducted (Netzsch, Bayern, Germany) at a heating rate of 10 C/min in static air with powder as a reference material. The densification kinetics, thermal expansion and coefficient of thermal expansion (CTE) were studied using high-temperature dilatometry (Netzsch). The flexural strength was measured at room temperature using a three-point bending strength universal testing machine (Model 4204, Instron Corp., Danvers, Mass.). Thermal shock resistance of the samples was studied by measuring the retained flexural strength after multiple water quenching cycles of 1000 C. The samples were heated to 1000 C in an electric furnace with a soaking time of 20 min, were quenched in water for 5 min and suddenly reheated at 1000 C for 20 min. The retained flexural strength was measured after 40 heating and cooling cycles. Three samples were taken to determine the retained flexural strength and average values were recorded. Raw-Materials Characterization The particle-size distribution of the starting materials in terms of D 10, D 50 and D 90 were determined (Table 3). All the starting materials were nanosized or micronsized. The D 50 of the was ~0.950 µm and that of and kaolin was 4.47 and 4.76 µm, respectively. The specific surface areas were 1.18 m 2 /g for, m 2 /g for and m 2 /g for kaolin. Chemical analyses of the raw materials were determined (Table 1). The SiO 2 content of the kaolin batch was higher than the stoichiometric amount for mullitization. Excess SiO 2 was present as quartz. 7 was added to reach the stoichiometric mullite composition. The used in this investigation was calcined that consisted only of corundum (α- ). Differential Thermal Analysis A DTA curve was recorded during heating of the Tieh kaolin (Fig. 1). The curve reveals an endothermic peak at ~110 C because of vaporization of absorbed water and at C because of loss of structural water (hygroscopic) with the formation of metakaolinite. The prominent exothermic peak at ~980 C corresponds to the formation of mullite spinel from the metakaolin. A much smaller exothermic peak occurs at ~1150 C, which is associated with 3:2 mullite formation. 8 A TGA curve for the kaolin also was recorded (Fig. 1). The clay sample lost ~12 wt% in two temperature regions. The first loss in weight (at ~110 C) is due to loss of surface water. The second loss in weight (at ~550 C) is due to loss of water of crystallization (OH groups attached to aluminum and silicon). 9 Both TGA changes are indicated by endothermic changes in the DTA curves. The absence of further change in weight at higher temperatures, after the dehydroxylation is complete, confirms that the exothermic change at ~980 C in the DTA of the sample is due only to phase change. Table 2 Batch Design of Composites Sample Kaolin code (wt%) (wt%) (wt%) MZ MZ MZ MZ MZ Table 3 Starting Material PSD Densification Variations of the bulk density of the mullite composites sintered Starting material D 10 (µm) D 50 (µm) D 90 (µm) at 1600 C with content were plotted (Fig. 2). In all cases, the bulk density increases with increase of content. The highest bulk density (2.84 g/cm ) is reached with an apparent porosity of ~5% (batch MZ15 sintered at 1600 C). The bulk density at Kaolin C is dependant on the content. It increases from 2.67 g/cm 3 for MZ5 to 2.84 g/cm 3 for MZ15. This may be due to the higher specific gravity of American Ceramic Society Bulletin, Vol. 85, No. 11

4 It is believed that the relatively lower densification with higher contents may result from the excessive thermal expansion mismatches between and mullite matrix. This may originate internal cracks and weaken the matrix, which results in lower densification. In addition, the higher the content, the higher the viscosity of the formed glassy phases and, consequently, the lower the particle diffusion and rearrangement. It has been reported 10 that, in the SiO 2 system, at least two routes for mullite formation are recognized. First, mullite is formed as a layer in the SiO 2 contact zone by interdiffusion of Al 3+ into the silicate glassy phase. This is a slow process because of slow aluminum-ion and silicon-ion diffusion. Second, the silicate-rich glassy phase crystallizes. The dissolution of in this glass is considered as the rate-controlling kinetics of mullite. The concentration of in this glass increases until the critical nucleation concentration (CNC) of mullite is reached, at which crystallization of mullite proceeds rapidly at higher temperatures. The rapid formation of mullite traps the pores. Therefore, closed pores are formed, which inhibit further densification. Table 4 Bending Strengths and CTE Values Bending strength Bending strength Coefficient before after of thermal thermal shock thermal shock expansion Batch (MPa) (MPa) ( 10 6 / C) MZ MZ MZ MZ MZ It has been concluded that the true density of the sintered mullite composites depends on the nature and amount of the various phases formed during sintering. Formation of glassy phase in the sintered masses can occur in two ways. The first is caused by impurities associated with the starting material. These impurities may form liquid phase at high temperature The second is caused by a complex interaction of the additives in the SiO 2 system, which leads to formation of several eutectic and lowmelting-point phases. 11 The densification occurs mostly through solid-state sintering, with small amounts of glass formation caused by the presence of impurities in the raw materials. 12 Phase Analysis and Microstructure XRD patterns of the samples fired at 1600 C were collected (Fig. 3). They show that the major crystalline phase is mullite. A small amount of monoclinic- (m-) is detected. These results are in agreement with an earlier report that reveals that the added can be entrapped during the densification process by interaction with the glassy phase at the sintering temperature by formation of an amorphous SiO 2 -rich zircon phase. The zircon mixture has a higher density than the mullite mixture. Therefore, the reaction between zircon and to form mullite and is accompanied by a volume expansion, which may result in some stresses in the fired samples. According to Le Chatelier s principle, such stresses may inhibit the reaction between zircon and. This allows a small amount of to remain in the sample, which is evidenced by a weak peak in the XRD pattern. 13 It is concluded 14 that the particles are intragranular in mullite composites fabricated using reaction sintering of zircon powder mixtures. Thus, few t- particles are detected using XRD, and they remain below the critical size for transformation because of the increase of diffusion distances by an increased mullite layer. Also, the observation of a limited phase using XRD may be due to TiO 2 impurities in the clay raw materials. In a similar work, the addition of TiO 2 results in the decrease of t- concentration in the ceramic body when the sintering temperature is increased. This possibly can be attributed to the decrease of the t m transformation temperature by the formation of a solid solution of TiO 2 in t-. 15 Scanning electron microscopy (SEM) photographs were made of samples MZ0, MZ5, MZ10, MZ15 and MZ20 fired at 1600 C (Fig. 4). Because densification mostly occurs through solid-solution sintering, the mullite grains formed are mostly equiaxed. The presence of slightly decreases mullite grain growth and decreases the amount of glassy phase. particles mostly occupy the intergranular position within the mullite matrix. Small amounts of finer particles also are observed in the intragranular position. Prochazka et al. 16 have concluded that enters into the mullite structure by solid-solution formation. Elemental analysis shows that the /SiO 2 molar ratio of mullite formed in composites is , which indicates the nonstoichiometric nature of the mullite, and 1.7 wt% enters in the mullite structure by solid-solution formation. American Ceramic Society Bulletin, Vol. 85, No

5 The microstructure of sample MZ0 (Fig. 4(a)) shows a homogenous microstructure of mullite matrix with some amorphous glassy phase located at the grain boundaries. It is concluded 2 that addition of enhances the mullitization with a corresponding decrease in density from 3.79 to 3.75 g/cm 3. 2 Sample MZ5 has been characterized using SEM by rounded grains uniformly distributed in a mullite matrix. The intergranular grains act as a pinning agent to decrease mullite grain growth. Micrographs of samples MZ10 and MZ15 reveal important coarsening after firing at 1600 C. A photomicrograph of sample MZ15 shows the dense nature of the microstructure. It is characterized by the crosslinked structure of mullite grains and grains. There are two types of : intergranular (rounded to semirounded in shape and located between mullite grains); and intragranular (rounded in shape and smaller and located within the mullite grains). The growth of intergranular occurs by grainboundary mass transport of Zr 4+ ions, whereas grain growth of intragranular occurs by diffusion of Zr 4+ ions through the mullite lattice. The difference in rate of Zr 4+ -ion diffusion is responsible for the difference in grain sizes. 17 A SEM photograph of MZ20 (Fig. 4(e)) shows large porosity and voids with glassy-phase formation. The particles are intergranular and intragranular, and the mullite grains have lost their columnar structure, as compared with MZ15. This may have been caused by the effect of pining of particles to decrease mullite formation. The porous structure of MZ20 agrees with previous work. This may be caused by the impurities or the rapid formation of mullite via a crystallization process at higher temperature. The rapid mullitization traps the pores. Therefore, many closed pores are formed, which inhibit further densification. 18 These results are similar to previous findings that concern thermodynamic analyses 19 where the reaction of Fe 2 to Fe 3 O 4 can take place to liberate O 2 -gas as the temperature reaches ~1500 C. Also, the microstructure of the fracture surface of the composites reveals pores and voids that affect density and strength. 20 Flexural Strength Flexural strength at room temperature has been determined (Fig. 5). -containing composites fired at 1600 C show a gradual increase in the bending strength compared with -free composites because of lower bulk density. addition increases the flexural strength of the sample sintered at 1600 C. The initial decrease in strength of MZ10 may be related to the elastic bond relaxation with the sintered temperature. 21 There are two reasons for final strength enhancement at higher amounts of addition. First is the glassy phase, which can lead to the healing of critical flaws or increase the apparent toughness during the sintering process. Second is the strengthening of the grain-boundary mechanism produced by a continuous solid solution at the grain boundary between and mullite. The solid solution enhances grain-boundary mass transport and, consequently, the sintering rate. 22 Thermal Shock Resistance Measurements of the bending strength values before and after thermal shock in addition to the thermal expansion values of the sample sintered at 1600 C were used to study the thermal shock behavior. Threepoint flexural strength values were observed after 40 thermal shock quenching cycles into water ( T = 1000 C) (Table 4). In all cases, there is gradual decrease in strength values. The strength degradation of the samples after 40 cycles thermal shock is due to thermal fatigue caused by subcritical crack growth. Presence of may have improved the thermal shock resistance of sample MZ20 because of the decrease of viscosity of the formed glassy phases in the sintered samples. 23,24 All the composites show a linear thermal expansion to 1500 C. There are slight differences in their CTE values, which increase with an increase of the content. About the Authors S.H. Kenawy, M. Awaad and H. Awad are research staff members of the National Research Centre, Cairo, Egypt American Ceramic Society Bulletin, Vol. 85, No. 11

6 References 1 S.-K. Zhao, Y. Huang, C.-A. Wang, X.-X. Huang and H.-K. Guo, Mullite Formation from Reaction Sintering of ZrSiO 4 /α- Mixtures, Mater. Lett., 57, (2003). 2 F. Temoche, L.B. Garrido and E.F. Aglietti, Processing of Mullite Zirconia Grains for Slip Cast Ceramics, Ceram. Int., 31, (2005). 3 C. Aksel and F. Komicezny, Mechanical Properties and Thermal Shock Behaviour of PSR333 Alumina Mullite Zirconia Refractory Materials, Glass Int., 1, (2001). 4 H.M. Jang, S.M. Cho and K.T. Kim, Alumina Mullite Zirconia Composites: Part II, Microstructural Development and Toughening, J. Mater. Sci., 2, (1997). 5 L.B. Garrido and E.F. Aglietti, Pressure Filtration and Slip Casting of Mixed Alumina Zircon Suspensions, J. Euro. Ceram. Soc., 21, (2001). 6 N. Claussen and J. Jahn, Mechanical Properties of Sintered In-Situ Reacted Mullite Zirconia Composites, J. Am. Ceram. Soc., 63 [3 4] (1980). 7 C.Y. Chen, G.S. Lan and W.H. Tuan, Preparation of Mullite by the Reaction Sintering of Kaolinite and Alumina, J. Euro. Ceram. Soc., 20, (2000). 8 V. Viswabaskaran, F.D. Gnanam and M. Balasubramanian, Mulltization Behaviour of Calcined Clay Alumina Mixtures, Ceram. Int., 29, (2003). 9 O. Castelein, B. Soulestin, J.P. Bonnet and P. Blanchart, The Influence of Heating Rate on the Thermal Behaviour and Mullite Formation from a Kaolin Raw Material, Ceram. Int., 27, (2001). 10 S. Maitra, S. Pal, S. Nath, A. Pandey and R. Lodha, Role of MgO and Cr 2 Additives on the Properties of Zirconia Mullite Composites, Ceram. Int., 28, (2002). 11 E. Bischoff and M. Ruhle, Thin Boundaries in the Particles Confined in a Mullite Matrix, J. Am. Ceram. Soc., 66, (1983). 12 M. Ruhle, N. Claussen and A.H. Heuer, Transformation and Microcraking as Complementary Processes in Zirconia- Toughened Alumina, J. Am. Ceram. Soc., 69, (1986). 13 J.H. She, H. Schneider, T. Inoue, M. Suzuki, S. Sodeoka and K. Ueno, Fabrication of Low-Shrinkage Reaction- Bonded Alumina Mullite Composites, Mater. Chem. Phys., 68, (2001). 14 M.G.M.U. Ismail, Z. Nakai and S. Somiya, Properties of Zirconia-Toughened Mullite Synthesized by Sol Gel Method ; pp in Advances in Ceramics, Vol. 24, T. Ebadzadeh and E. Ghasemi, Effect of TiO 2 Addition on the Stability of t- in Mullite Composites Prepared from Various Starting Materials, Ceram. Int., 28, (2002). 16 S. Prochazka, J.S. Wallace and N. Claussen, Microstructure of Sintered Mullite Zirconia Composites, J. Am. Ceram. Soc., 66, C-125 C-127 (1983). 17 K. Das, S.K. Das, B. Mukherjee and G. Barnejee, Microstructural and Mechanical Properties of Reaction-Sintered Mullite Zirconia Composites with Magnesia as Additive, Interceram, 46 [5] (1998). 18 S. Zhao, X. Huang and J. Guo, The Effect of Mullite Seeding on Reaction-Sintered Mullite Zirconia Multiphase Ceramic, Mater. Sci. Lett., 19, (2000). 19 C.Y. Chen and W.H. Tuan, The Processing of Kaolin Powder Compact, Ceram. Int., 27, (2001). 20 D.R. Gaskell, Introduction to Metallurgical Thermodynamics, 2nd ed.; p Scripta Publishing, M.I. Osendi and C. Baudin, Mechanical Properties of Mullite Materials, J. Euro. Ceram. Soc., 96, (1996). 22 J.S. Moya and M.I. Osendi, Effect of (SS) in Mullite on the Sintering and Mechanical Properties of Mullite/ Composites, Mater. Sci. Lett., 2, (1983). 23 J. Gebauer, D.A. Krohn and D.P.H. Hasselman, Thermal-Stress Fracture of a Thermomechanically Strengthened Aluminosilicate Ceramic, J. Am. Ceram. Soc., 55, (1972). 24 H.P. Kirchner, R.E. Walker and R.M. Gruver, Strengthening Alumina by Quenching in Various Media, J. Appl. Phys., 42, (1971). American Ceramic Society Bulletin, Vol. 85, No

7 Apparent porosity (%) Bulk density (g/cm 3 ) Zirconia content (wt%) Fig. 2 Densification parameters (( ) apparent porosity and ( ) bulk density) of the fired samples as a function of content. Intensity (a.u.) 2θ (deg) Fig. 3 XRD patterns of ceramic composite samples fired at 1600 C (( ) mullite, ( ) t- and ( ) unknown).

8 Flexural strength (MPa) Bulk density (g/cm 3 ) (wt%) Fig. 5 Relationship between content and ( ) flexural strength and ( ) bulk density values of composite fired at 1600 C.