Corrosion and Oxidation Behaviour of Β-Sialon Ceramics Via Different Processing Route

Journal of the Ceramic Society of Japan 117 [4] 482-488 2009 Paper

Corrosion and oxidation behaviour of β-SiAlON ceramics via different processing route

T. PLACHKÝ,† J. KRˇ ESTˇ AN,* M. KORENKO, V. MEDRI,** Z. LENCˇ ÉŠ and P. ŠAJGALÍK

Institute of Inorganic Chemistry, Slovak Academy of Sciences, Dúbravská cesta 9, SK-845 36 Bratislava 45, Slovak Republic *VOP-026 Sternberk, s. p., Division VTUO Brno, Veslarska 230, 637 00 Brno, Czech Republic **Institute of Science and Technology for Ceramics, National Research Council, Via Granarolo 64, 48018 Faenza, Italy

The corrosion resistance of β-SiAlON based advanced ceramic materials in molten aluminum, equimolar NaCl–KCl melt and also in the melted mixture of NaF and AlF3 (molar ratio 1 : 1.163) were investigated. Two kinds of β-SiAlONs were tested; β- SiAlON prepared from powder precursors produced by carbothermal reduction and nitridation (CRN) of pyrophyllite and β- SiAlON made from commercial powders (AlN, Al2O3 and Si3N4). Corrosion tests were realized at 760°C for 7, 36 and 72 h under nitrogen atmosphere by dip-finger test method. All types of β-SiAlONs have great corrosion resistance against molten aluminum and NaCl–KCl melt at these conditions. The results of corrosion test were different for the β-SiAlONs in molten fluorides. While the corroded zone was only about 100 μ m deep after 72 h corrosion test in reference β-SiAlON prepared from synthetic powders (SiAlON-R), in sample prepared from CRN pyrophyllite (SiAlON-P) it was almost 230 μ m thick. Addition- ally, the oxidation resistance of β-SiAlONs was tested. The results showed that both SiAlON-P and reference SiAlON-R have comparable good oxidation resistance. ©2009 The Ceramic Society of Japan. All rights reserved.

Key-words : β-SiAlON, Carbothermal reduction and nitridation, Corrosion, Oxidation, Aluminum industry

[Received April 9, 2008; Accepted January 15, 2009]

are subsequently dissolved and oxidized by the formed Si–Li– 1. Introduction Al–Na–O–F melt. Corrosion in the Al–Li metal zone was much The application of advanced ceramics increases also in the less than in the flux and gas zones. The results are in agreement metallurgical industry and consequently the knowledge of their with previous studies in terms of the resistance of SiAlON to the corrosion behavior is essential.1) β-SiAlON ceramics with gen- attack by aluminum. However, formation of protective AlN layer eral formula Si6–zAlzOzN8–z, where 0 < z < 4.2, are frequently at the interface was not observed. used materials in the industry owing to their good mechanical The good performance of β-SiAlONs in foundry requires their properties, corrosion resistance and reasonable price. The use of cost effective high-scale production. β-SiAlON ceramics are β-SiAlONs in foundries significantly increased in recent years. usually prepared by reactive sintering of appropriate mixtures of SiAlONs are resistant to attack by several molten metals, what Si3N4, Al2O3 and AlN powders. Because the nitride powders are can be utilized also in aluminum foundry for enhancement of relatively expensive, many attempts have been made to convert durability and for reducing the contamination of the aluminum natural aluminosilicates to β-SiAlONs by carbothermal melt and its alloys. reduction and nitridation (CRN).8)–14) Aluminosilicate-derived The corrosion performance of Si3N4 in molten aluminum was SiAlONs are considered as a low cost alternative to those studied by Schwabe et al.2) Formation of protective AlN or AlN- available on the market. These β-SiAlONs are frequently used in polytype layer was observed at the interface between dense Si3N4 metallurgy, especially in aluminum foundry, e.g. by Kyocera and aluminum metal, which improved the corrosion resistance of Corp., Hitachi Metals Ltd., Kubota Corp., etc. However, there is studied Si3N4 materials. However, it was also shown that open still a lack of data on the corrosion resistance of β-SiAlONs porosity degrades the corrosion resistance of Si3N4 ceramics. depending on the used aluminosilicate precursor used for their Mouradoff et al. investigated the evolution of contact angle of production. molten aluminum on two types of silicon nitride as a function of From that reason in this work two types of β-SiAlONs were time and temperature.3),4) They confirmed the formation of inert prepared: first one sintered from precursor made by carbothermal AlN film on the surface of silicon nitride at temperatures above nitridation of pyrophyllite, the second one was prepared as ref- 1000°C. High corrosion resistance of Si3N4-based materials erence material from commercial powders of Si3N4, Al2O3 and against liquid aluminum was reported also by Oliveira et al.5) and AlN. Their corrosion resistance against molten aluminum and Santos.6) relative fluoride and chloride fluxes used in the aluminum Dower and Coley studied the corrosion of low z β-SiAlON by metallurgy has been studied. molten aluminum-lithium alloy and NaF–NaCl flux.7) Their Primary aim of this study is to determine the extent of corro- results showed that SiAlON ceramics are strongly attacked by sion of aluminosilicate derived β-SiAlONs in molten aluminum, the fluoride flux along the grain boundaries and β-SiAlON grains equimolar mixture of NaCl and KCl melt and also in the melted mixture of NaF and AlF3. Because the β-SiAlON products used † Corresponding author: T. Plachky; E-mail: [email protected] in metallurgy are working in hot oxidizing environments before

coming into contact with molten media, also their oxidation During the CRN process the impurities like Fe2O3 were trans- resistance is an important factor influencing the extent of their formed to FexSiy, TiO2 was reduced and nitrided to TiN, and CaO application. From that reason the oxidation of β-SiAlONs was dissolved in the silicate based glassy phase.14) prepared in this work was studied in both isothermal and non- The powder products of CRN were pressed into pellets in a isothermal conditions. steel die at 100 MPa. The compacted samples were hot-pressed in a graphite die at 1750°C for 2 h in 0.1 MPa of nitrogen and 2. Experimental procedure mechanical load of 30 MPa (samples SiAlON-P). Reference Two types of β-SiAlON ceramic materials were prepared from mixture R was also hot-pressed under the same conditions for two different starting powders. Mixture labeled as P in Table 1 comparison (samples SiAlON-R). XRD and SEM methods were was made from oxynitride precursors produced by carbothermal used to analyze the samples after high temperature treatment. reduction and nitridation of pyrophyllite (grade P–I, Envigeo The final z-value was 3.7 for both types of β-SiAlONs. The XRD Ltd., Slovakia) with appropriate amount of carbon black patterns of dense SiAlON-P and R samples (Fig. 1) show β- (Stickstoffwerk Piesteritz, Germany), β-Si3N4 (Tschernogolovka, SiAlON as the major phase in both cases. The presence of minor Russia) and α-Al2O3 (Martoxide PS–6, Martinswerk, Germany). phases Al2O3, SiC, FexSiy and TiN was observed in pyrophyllite The chemical composition of pyrophyllite is given in the derived SiAlON-P. Table 2. The major and not negligible impurities are Fe2O3, TiO2 The dense bars of both SiAlON-P and SiAlON-R samples and CaO. The pyrophyllite-based starting powder was attritor- were investigated by static corrosion test method (Fig. 2). Sam- milled for 0.5 h in isopropanol, dried and sieved through 70 μ m ples were placed into alsint crucibles and dusted with aluminum screen. powder (Hichem Comp. Ltd., Czech Republic). Equimolar mix- The reference β-SiAlON samples labeled as SiAlON-R were ture of NaCl and KCl was placed on the top of the batch. Other prepared from mixture R (Table 1) consisting of commercial three cuboids were placed into special crucible with a mixture of powders: AlN (grade C, H.C. Starck, Germany), above NaF and AlF3 (molar ratio 1:1.163 corresponding to the second 15) mentioned α-Al2O3, and α-Si3N4 (Siconide-P, Permascand eutectics of binary NaF–AlF3 phase diagram). Crucibles were Ceramic, Sweden). The starting compositions of both powder inserted into a vertical resistance furnace and annealed at 760°C mixtures (Table 1) were adjusted to obtain β-SiAlON with z = 4 for 7, 36 and 72 h in N2 atmosphere. The polished cross-sections in Si6–zAlzOzN8–z after sintering. of heat-treated samples were analyzed by SEM (Zeiss EVO 40) The homogenized pyrophyllite-based powder mixture P was equipped with EDX analyzer (Bruker AXS, Quantax 400 with poured into graphite reactor, which was inserted into rotary silicon drift detector). furnace. Nitrogen gas was introduced at 600°C through a diffuser The isothermal oxidation tests of SiAlON-P and SiAlON-R into the reactor. The nitrogen flow (100 cm3/min) was controlled samples were carried out at 800°C, 1000°C, 1200°C and 1400°C by means of calibrated flow-meter. The final temperature of CRN for 20 h with heating rate 30°C/min. In the case of the non was 1510°C with 8 h dwell for pyrophyllite batches.14) The phase isothermal oxidation the furnace was heated to maximum tem- composition and morphology of CRN products were analyzed by perature 1450°C with the fixed heating rate 2°C/min and air flow XRD (STOE STADI-P, CoKα1, Germany) and SEM (Zeiss 30 ml/min. The weight gain was continuously recorded with a EVO40, Germany) methods. The major phase was β-SiAlON. thermogravimetric apparatus (STA 449 C Jupiter®-Simultaneous TG–DSC, NETZSCH Gerätebau GmbH, Germany). The static long-term oxidation tests were carried out also for 100 h at tem- Table 1. Composition of Starting Powder Mixtures in mass% peratures 1000°C, 1200°C and 1400°C in air. After oxidation the specimen’s weight changes were evaluated. Bending strength of Sample P–I Carbon β-Si3N4 α-Al2O3 AlN α-Si3N4 P 53.5 25.8 5.2 15.5 R 47.9 32.9 19.2

Table 2. Chemical Composition of Pyrophyllite P–I in mass% (According to Envigeo, Ltd.)

SiO Al O Fe O TiO CaO MgO Na O K O LOI. Label 2 2 3 2 3 2 2 2 Σmass% (mass%)(mass%)(mass%)(mass%)(mass%)(mass%)(mass%)(mass%)(mass%)

P–I 78.23 15.85 1.65 0.59 0.23 0.01 0.02 0.03 3.11 99.72

Fig. 1. XRD patterns of β-SiAlON made from pyrophyllite (P) and commercial powders (R) after hot pressing at 1750°C/2 h/30 MPa. Fig. 2. Schematic of the corrosion experiments.

483 JCS-Japan Plachký et al.: Corrosion and oxidation behaviour of β-SiAlON ceramics via different processing route

Fig. 3. SEM micrographs of β-SiAlON samples after corrosion test at 760°C for 72 h in NaCl–KCl flux.

by molten aluminum was observed at 760°C in N2 atmosphere. Owing to the low wetting of SiAlON by molten aluminum, the flux of chlorides penetrated during rotary tests to the certain depth between the β-SiAlON test-bars and metal, as it is sche- matically illustrated in Fig. 4. In order to avoid any misinterpre- tation of corrosion results, the polished cross-sections for EDX analysis were prepared from the zone, which was not affected by the presence of chlorides. There is no evidence of the corrosion of the surface of both tested SiAlON-R and SiAlON-P materials by molten aluminum even after 72 h treatment at 760°C, as it is shown in Fig. 5. Line scans of aluminum, silicon, oxygen and nitrogen taken from the SiAlON–aluminum interface confirmed this observation (Fig. 5). The formation of protective AlN layer on β-SiAlON surface was not observed at the conditions of experiment. Contrary to pure Si3N4, where AlN or AlN-polytype layer forms on the interface between dense Si3N4 and molten aluminum, β-SiAlON suppresses this reaction. This is related to the Fig. 4 . Schematic of the penetration of chlorides into the melted alumi- high value of z-parameter in both pyrophyllite-derived and refer- num zone. ence β-SiAlONs (z = 3.7). High z-value implies the high level of aluminum substitution in β-SiAlONs resulting in low driving force for the reaction between molten aluminum and β-SiAlON. SiAlONs (P and R) was measured after long-term oxidation at 1000°C by 4–point bending test at room temperature. The strength 3.3 Interactions between β-SiAlON and NaF–AlF3 of sample SiAlON-P was tested also after treatment at 25°C, flux 1200°C and 1400°C. The surface and cross-section of ceramics of The thickness of corrosion layer of tested β-SiAlONs in oxidized ceramics were investigated by SEM–EDX method. molten fluorides is rather different. The pyrophyllite-derived SiAlON-P has more than two times thicker corroded zone 3. Results and discussion compared to reference SiAlON-R after 7, 36 and 72 h at 760°C. 3.1 Interactions between β-SiAlON and NaCl-KCl Similar behavior of these materials was observed after corrosion flux tests in molten steel.17) Both types of the prepared β-SiAlON materials did not exhibit Fluoride melt penetrated the pyrophyllite-derived SiAlON-P any signs of corrosion after testing at 760°C for 7, 36 and 72 h through grain boundaries to the depth of 25 μ m after 7 h treat- in NaCl–KCl flux, as it is shown in Fig. 3. The interface between ment. Thickness of the attacked zone in the reference SiAlON- ceramics and flux, part of which was solved out during cutting R is only about 10 μ m. The polished cross section of the and polishing in water solutions, is continuous and clear without corroded zone of samples is similar to the microstructure of any visible damage of the surface or change of microstructure of liquid phase sintered ceramics after chemical etching. SiAlON ceramics. This observation was confirmed by EDX analysis. matrix grains with etched out grain boundaries are clearly visible Although the surface of β-SiAlON has been wetted by molten in Fig. 6. These results indicate that the oxygen-rich grain fluxes of chlorides, no interaction was observed. The high cor- boundaries are attacked first by the fluoride melt. The observa- rosion resistance of aluminosilicate-based β-SiAlONs in the melt tion is consistent with the results of Jorge et al.,18) who showed of chlorides is consistent with the data reported by Takeuchi et that silicon oxynitride Si2N2O and alumina have lower resistance al.16) for silicon nitride immersed in NaCl–KCl molten salt at to cryolite bath compared with that of silicon nitride. Dower et 7) similar temperatures, but under Cl2–O2 atmosphere for 480 h. al. also observed that the silica-based grain boundaries are attacked first by the fluoride during the corrosion tests of 3.2 Interactions between β-SiAlON and molten SiAlONs in NaF–NaCl system. Afterwards, the fluorides are aluminum penetrating deeper into the ceramics. Negligible or no-wetting of both types of prepared β-SiAlONs The thickness of corroded zone increased to about 80 μ m in

484 Journal of the Ceramic Society of Japan 117 [4] 482-488 2009 JCS-Japan

Fig. 5. Linescan analysis of the cross section of SiAlON-P made from pyrophyllite (left) and reference SiAlON-R (right) after test in molten aluminum at 760°C for 72 h.

composition to the previous case. However, there is a difference in the thickness of corrosion zones. The depth of the first region with Al2O3 and SiAlON grains embedded in NaF–AlF3 matrix is 30 μ m (zone 1 in the Fig. 7(b)). The thickness of the zone 2 is around 70 μ m and contains β-SiAlON grains without grain boundary phase, which was etched out by the fluorides. From these observations it can be concluded that the corrosion mechanisms in pyrophyllite-derived SiAlON-P and reference SiAlON-R is similar. However, the kinetics of the corrosion is different, which is twice higher for the SiAlON-P owing to the presence of different impurities introduced by the starting pyrophyllite powder. The backscattered SEM image of the hot pressed SiAlON-P sample (Fig. 8) prior to the corrosion test shows white grains in the whole bulk of ceramics. EDX analysis Fig. 6. Micrograph (BSE) of reference SiAlON-R after corrosion at of these grains confirmed the presence of iron and silicon (point ° 760 C for 36 h in fluoride melt. The corroded grain boundaries are visible. 2 in the Fig. 8) or titanium with higher amount of nitrogen (point 3 in the Fig. 8). FexSiy and TiN were formed during the reduction and nitridation of pyrophyllite from the impurities of Fe2O3 and 14) samples SiAlON-P after 36 h treatment. SEM/EDX investiga- TiO2. All these impurities decrease the resistance of SiAlON tions showed that the corroded zone is divided into two regions. against fluorides. Moreover, SiAlON-P contained some 14),17) The first area consists of three different phases to the depth of unreacted alumina. These remaining Al2O3 grains can be 15 μ m from the original interface: the continuous dark grey dissolved in the fluoride melt, because the solubility of Al2O3 in 19) matrix of fluoride flux contains small white isolated grains with Na3AlF6–AlF3 eutectic system at 684°C is 3.2 mass%. The dis- high iron content and light grey SiAlON grains, which were solved alumina grains can facilitate the penetration of the fluo- partly corroded. The second region contains β-SiAlON grains rides into the ceramic bulk. In the final step β-SiAlON grains are with fluoride flux located along grain boundaries up to the depth decomposed in fluoride environment to Al2O3 grains and to other of 80 μ m. Compared to the pyrophyllite-derived SiAlON-P the corrosion products like gaseous SiF4 and N2. The in situ formed thickness of the corroded region in reference sample SiAlON-R alumina can further dissolve in fluoride melt up to its solubility is only about 40 μ m thick. The attacked β-SiAlON grains are limit. located up to the depth of 5 μ m from the surface. Undamaged grains of SiAlON with corroded grain boundary phase can be 3.4 Oxidation resistance of β-SiAlONs (P and R found beneath this layer. In the depth of 40 μ m only the original samples) dense β-SiAlON is visible. Because the parts of ceramic tools used in aluminum industry The corroded zone of SiAlON-P in fluoride melt is 230 μ m are exposed to high temperatures in oxidizing atmosphere, the thick after 72 h of the treatment (Fig. 7(a)). The top 130 μ m thick oxidation resistance of β-SiAlONs was also investigated. The layer (zone 1) is almost disrupted. Dominant phase in this region isothermal thermogravimetry tests lasting for 20 h did not show is the fluoride flux (points 1, 4 and 5 in Fig. 7(a)), in which Al2O3 any weight change up to 1200°C in both tested SiAlON-P and grains can be observed (point 3 in Fig. 7(a)) and some remaining SiAlON-R materials. Slight increase of the weight of oxidized β-SiAlON grains. Except corrosion, large pores were found in samples was observed at 1400°C, as it is shown in Fig. 8. The this layer formed by gaseous SiF4 and N2 reaction products. normalized thermogravimetric curves of both β-SiAlONs show Zone 2 consists of β-SiAlON (point 6 in Fig. 7(a)) and grain (Fig. 8) that the weight gain followed the parabolic law reported boundaries corroded by fluorides. Undamaged β-SiAlON with- also in the literature.20)–22) out the presence of fluorides is in the depth of 230 μ m from the On the other hand, the non-isothermal oxidation test showed original interface. The corrosion layer in the reference SiAlON- a rapid increase of oxidation rate above 1200°C, Fig. 9. This R is only 100 μ m deep, Fig. 7(b) and similarly to the SiAlON-P threshold of abrupt oxidation increase is about 150°C higher it can be divided into two different zones having comparable compared to the literature data obtained for β-SiAlON with

485 JCS-Japan Plachký et al.: Corrosion and oxidation behaviour of β-SiAlON ceramics via different processing route

Fig. 7. SEM backscattered images of the corrosion region of SiAlON-P (a) and SiAlON-R (b) after test in NaF–AlF3 flux at 760°C for 72 h. Elemental EDX analysis for particular spots is also attached.

Fig. 8. SEM backscattered image of SiAlON-P with three points analyzed by EDX: 1 - β-SiAlON; 2 - iron silicide (FexSiy) and SiAlON; 3 - TiN and SiAlON.

486 Journal of the Ceramic Society of Japan 117 [4] 482-488 2009 JCS-Japan

z = 3.12) The abrupt oxidation at 1200°C was confirmed also by this observation, Fig. 12. The strength of SiAlON-P sample long term static tests, when SiAlON-P and SiAlON-R samples decreased after oxidation at 1000°C for 100 h by ~18%, however were oxidized at different temperatures for 100 h (Fig. 10). The the majority of fracture origins were in the bulk of the tested sample weight gain normalized to the surface (Δm/S) is substan- bars. Substantially higher strength decrease of about 55% was tially changed at temperatures above 1200°C. observed for the sample oxidized at 1400°C for 100 h. The The parabolic character of the oxidation curve at the isother- fracture origin was located in the oxidized surface layer of this mal test (Fig. 8) indicates the formation of superficial oxidized sample in all tested bars. layer, which protects the ceramics against further oxidation. 4. Conclusions SEM/EDX investigation of the surface after 20 h showed the formation of mullite crystals at 1400°C and the presence of Both tested β-SiAlON based ceramics, SiAlON-P prepared aluminosilicate glassy phase from 1200°C. EDX analysis of the from pyrophyllite by carbothermal reduction and nitridation corroded layer after 100 h oxidation at 1400°C confirms the for- (CRN) process and the second reference SiAlON-R prepared mation of mullite on the surface, Fig. 11. The thickness of the from commercial powders exhibit good corrosion resistance oxidized layer was around 30 μ m at this temperature, while against molten chloride salt (NaCl–KCl) and liquid aluminum. almost no corrosion layer was found after oxidation bellow No evidence of interaction between β-SiAlONs and chloride or 1000°C. Bending strength of corroded samples corresponds to aluminum melt was observed at 760°C after 72 h treatment. In terms of application of β-SiAlON made from precursors prepared by CRN of raw aluminosilicates (e.g. pyrophyllite), their good corrosion resistance in such environments is expected. In the case of tests in NaF–AlF3-based fluoride melt, the reference β-SiAlON-R showed better corrosion resistance compared to hydrosilicate derived SiAlON-P. The corroded zone of SiAlON-P sample is more then two times deeper compared to the reference SiAlON-R, owing to the presence of different impurities from the raw hydrosilicate. The average thickness of corrosion zone after 72 h at 760°C is 230 μ m in SiAlON-P and 100 μ m in SiAlON-R, respectively. However, the main corrosion mechanism is the same. Grain boundary phase is attacked first by the fluorides, subsequently the β-SiAlON grains react with

Fig. 9. Time dependence of weight gain normalized to the sample surface for SiAlON-P and SiAlON-R at 1400°C for 20 h.

Fig. 11. Temperature dependence of weight gain (W) and degradation Fig. 10. Temperature dependence of weight gain normalized to the of bending strength (σ ) after 100 h holding time at particular tempera- Δ sample surface m/S for SiAlON-P and SiAlON-R. tures.

Fig. 12. Micrograph of the cross-section of pyrophyllite derived SiAlON-P after long-term oxidation at 1400°C with elemental analysis of the oxidized layer.

487 JCS-Japan Plachký et al.: Corrosion and oxidation behaviour of β-SiAlON ceramics via different processing route

the surrounding fluoride flux to form Al2O3 and gaseous SiF4 and (1996). N2. Different rate of corrosion in the two kinds of β-SiAlONs, 8) I. Higgins and A. Hendry, Br. Ceram. Trans. J., 85, 161–166 however with similar basic composition of z = 3.7, is caused by (1986). 9) K. J. D. MacKenzie, R. H. Meinhold, G. V. White, C. M. the presence of impurities and minor phases (FeSix, TiN, Al2O3) in the hydrosilicate-derived β-SiAlONs. Sheppard and B. J. Sherriff, J. Mater. Sci., 29, 2611–2619 (1994). The good oxidation resistance of β-SiAlON prepared from 10) J. Mukerji and S. Bandyopadhyay, Adv. Ceram. Mater., 3, pyrophyllite is comparable to the reference SiAlON-R. 369–373 (1988). β Hydrosilicate-derived -SiAlONs could be good and cheaper 11) Y. Sugahara, K. Kuroda and C. Kato, J. Mater. Sci., 23, 3572– candidates for application in aluminum industry, especially in 3577 (1988). processes, where direct contact with cryolite melt is avoided. 12) T. Bastow, S. G. Hardin and T. W. Turney, J. Mater. Sci., 26, According to the obtained results β-SiAlONs can be also used 1443–1453 (1991). during next processing steps, i.e. in foundry during casting and 13) G. Sucik, T. Kuffa and D. Hrsak, Metalurgija, 43, 93–95 manipulation with molten aluminum. (2004). 14) J. Krestan, P. Sajgalik and Z. Panek, J. Eur. Ceram. Soc., 24, Acknowledgments This work was partly supported by the 791–796 (2004). National Slovak Grant Agency VEGA, project No 2/7171/27, 15) A. Solheim and A. Sterten, Proceedings of the Ninth APVV-0171-06 and NANOSMART, Centre of Excellence, SAS. International Symposium on Light Metals, Trondheim, Norway, 225–234 (1997). 16) M. Takeuchi, T. Kato, K. Hanada, T. Koizumi and S. Aose, J. References Phys. Chem. Solids, 66, 521–525 (2005). 1) F. L. Riley, Key Eng. Mater., 113, 1–14 (1996). 17) J. Krestan, O. Pritula, L. Smrcok, P. Sajgalik, Z. Lences, A. 2) U. Schwabe, L. R. Wolff, F. J. J. Van Loo and G. Ziegler, J. Wannberg and F. Monteverde, J. Eur. Ceram. Soc., 27, 2137– Eur. Ceram. Soc., 9, 407–415 (1992). 2143 (2007). 3) L. Mouradoff, A. Lachau-Durand, J. Desmaison and J. C. 18) E. Jorge, O. Marguin, L. Moitrier and O. Citti, USPTO Appli- Labbe, J. Eur. Ceram. Soc., 13, 323–328 (1994). caton No. 20080234122 - Class: 501 92, (2008). 4) L. Mouradoff, P. Tristant, J. Desmaison, J. C. Labbe, M. 19) P. A. Foster, J. Am. Ceram. Soc., 58, 288–291 (1975). Desmaison-Brut and R. Rezakhanlou, Key Eng. Mater., 113, 20) J. Persson and M. Nygren, J. Eur. Ceram. Soc., 13, 467–484 177–186 (1996). (1994). 5) M. Oliveira, S Agathopoulos and J. M. F. Ferreira, J. Eur. 21) R. Ramesh, P. Byrne and M. J. Pomeroy, J. Mater. Process. Ceram. Soc., 25, 19–28 (2005). Tech., 56, 600–608 (1996). 6) C. Santos, S. Ribeiro, K. Strecker, D. Rodrigues Jr. and C. R. 22) X. M. Zhou, K. C. Chou and F. S. Li, J. Eur. Ceram. Soc., 28, M. Silva , J. Mater. Process. Tech., 184, 108–114 (2007). 1243–1249 (2008). 7) L. T. Dower and K. Coley, Key Eng. Mater., 113, 167–176

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