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Joining of ceramic oxides by liquid wetting and capillarity
Scripta Materialia 45 (2001) 759±766
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Joining of ceramic oxides by liquid wetting and capillarity Laura Esposito and Alida Bellosi* IRTEC CNR, Research Institute for Ceramics Technology, via Granarolo 64, 48018 Faenza RA, Italy
Received 5 February 2001; accepted 12 June 2001
Abstract A bonding technique based on the use of a glass powder interlayer is studied. Examples are presented for the systems formed by ZrO2 (3 mol%Y2 O3 ) and the composite 80Al2 O3 ±20ZrO2 (wt.%) and a silicate glass as interlayer. The microstructure and mechanical resistance of the joints are also reported. Ó 2001 Published by Elsevier Science Ltd. on behalf of Acta Materialia Inc. Keywords: Bonding; Ceramics; Oxides; Glasses; Interfaces
Introduction Polycrystalline oxide materials like Al2 O3 or ZrO2 and related composites have been widely investigated in the last decades in terms of microstructure [1±6] and thermomechanical properties [7±9]. Most of the characteristics of these oxides are related and governed by the microstructure, i.e. the mean grain size and distribution, the composition and thickness of the intergranular phase, etc. These features play a key role in determining the ®nal properties of the material and depend primarily on the speci®c cations segregation (from impurities and/or sintering aids) occurring during sintering and on the consequent interfacial force balance between the grains [10]. The interactions occurring at high temperature among the phases that come in contact are fundamental also to produce adhesion between two dierent materials. The wetting on ¯owing capability of the intergranular phase, is a key factor for example for the production of bonds through the use of a pressure applied on two ceramic pieces in contact [11±14]. As an alternative to pressure-assisted joining methods, a glass or glass±ceramic interlayer placed between two ceramic pieces can be used to promote adhesion [15±17]. *
Corresponding author. E-mail address: [email protected] (A. Bellosi).
1359-6462/01/$ - see front matter Ó 2001 Published by Elsevier Science Ltd. on behalf of Acta Materialia Inc. PII: S 1 3 5 9 - 6 4 6 2 ( 0 1 ) 0 1 0 9 2 - 2
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L. Esposito, A. Bellosi/Scripta Materialia 45 (2001) 759±766
No pressure is needed with this technique and, if the glass composition is properly selected, the new phases which may form at the interface are not deleterious for the mechanical resistance of the joint. The optimal glass composition is function of the ceramic material, in particular: ± the ceramic must be chemically and thermomechanically compatible with the glass; ± the glass must wet the ceramic; ± the glass must melt below the temperature at which the ceramic starts to degrade (i.e., below the sintering temperature); ± the glass must have a relatively low viscosity in order to ¯ow easily between the two ceramic parts. In the following, two systems exhibiting a dierent behavior during the bonding cycle are presented. The same Ca±Al±silicate glass is used as interlayer for the composite Al2 O3 ±ZrO2 and for ZrO2 (3 mol%Y2 O3 ). The penetration of the glass in the ceramic and the mobility and solubility of the ceramic grains in the glass drive the ®nal interfacial microstructure. Experimental The compositions and principal characteristics of the ceramic materials are reported in Table 1. The glass used as interlayer (prepared at the Polytechnic of Torino, Italy) has the composition SiO2 39, Al2 O3 30, CaO 31 (wt.%), which falls in the primary phase ®eld for anorthite CaO Al2 O3 2SiO2 [18], has a melting point of 1430°C, as resulted from the heating microscope analysis, a coecient of thermal expansion of 6:6 10 6 K 1 [19±21] and a density of 2.76 g/cm3 . The ceramic surfaces kept in contact during the bonding cycles were previously polished with diamond paste with two dierent grades of surface ®nishing: (i) 1 lm diamond paste was used when the microstructural evolution of the two surfaces in contact and the ®nal glass thickness were analyzed; (ii) a polishing level of 15 lm is adopted to evaluate the feasibility and resistance of joints through an experimental procedure which can be compatible with industrial processes. The glass amount required to obtain the desired thickness was calculated considering the surface area of the ceramic piece and the density of the glass. The glass powder was sieved below 25 lm, weighed, ultrasonically dispersed in ethanol and deposited on the polished and cleaned surface of one ceramic piece. The other ceramic part was placed on top of the deposited glass and the assembled sample was kept at 60°C for 24 h to remove completely the ethanol. Joints were obtained using a laboratory furnace in air, under the experimental conditions reported in Table 2. The interfacial microstructure of Table 1 Compositions, sintering conditions and some characteristics of the materials Material
Sintering cycle (°C h)
Density gr/cm3 (%T.D.)
a (10 6 /K)
r (MPa)
ZrO2 (3 mol% Y2 O3 ) Al2 O3 ±ZrO2 (3 mol% Y2 O3 )
1500 1 1600 1
6.04 (99.8) 4.22 (98.4)
10 8.5
600 20 430 30
a: thermal expansion coecient (20±1200°C), r: 4-pts R.T. ¯exural strength. Experimental procedure described elsewhere for Al2 O3 ±ZrO2 [9], ZrO2 [22,23]. The composition of the composite is 80Al2 O3 ±20ZrO2 (wt.%).
L. Esposito, A. Bellosi/Scripta Materialia 45 (2001) 759±766
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Table 2 Bonded systems (ZX for ZrO2 =glass joining tests, AZX for Al2 O3 ±ZrO2 =glass joining tests); bonding conditions and 3-pts R.T. ¯exural strength (r) values of the joints. A Anorthite CaO Al2 O3 2SiO2 Sample
Cycle (°C min)
Starting glass thickness (lm)
Polishing level (lm)
Glass penetration (lm)
New crystal phases in glass
Final glass thickness (lm)
r (MPa)
Z1 Z2 Z3 Z4 Z5 Z6
1500 30 1500 30 1500 30 1450 30 1450 15 1450 15
± 20 50 50 50 30
1 1 1 1 15 15
450 200 350 200 200 gl.
ZrO2 None None None None None
± No glass No glass No glass 10 0±3
± ± 173 88 99 56 47 29 ±
AZ1 AZ2 AZ3 AZ4 AZ5 AZ6 AZ7
1500 30 1450 30 1450 30 1450 30 1450 15 1450 15 1450 15
± 50 20 50 20 30 50
1 1 1 15 1 15 15
ZrO2 , Al2 O3 ZrO2 , A None ZrO2 , A None Partiala Partiala
± 50 20 40 10 25±30 25±30
± 148 39 152 67 191 115 156 82 ± ±
a
Crystallization of ZrO2 and anorthite occurred in the central part of sample, not at the edges.
the joints was analyzed by optical microscopy (OM), scanning electron microscopy (SEM) and energy dispersive spectroscopy (EDX). The fracture strength of the bonds was evaluated in 3-point bending on bars 14:0 2:0 1:5 mm3 on a jig of 11 mm span with a crosshead speed of 0.5 mm/min and calculated with the usual bending formula. Results ZrO2 =CaO±Al2 O3 ±SiO2 glass system A wetting experiment (Z1) and various joining cycles (Z2±Z6) were performed with the zirconia/glass system (Table 2). In the wetting experiment the glass wets the ceramic with a contact angle h < 10° and penetrates extensively through the grain boundaries up to a distance from the interface of about 400±500 lm (Fig. 1a). The composition of the residual glass at the interface resulting from the EDX semiquantitative analysis is Al2 O3 30, SiO2 33, CaO 27, ZrO2 8, Y2 O3 2 (wt.%). The content of Ca and Si is less than in the original glass, whereas the content of Zr and Y is higher. These elements diused from the ceramic as a consequence of the partial dissolution of zirconia grains into the glass. In addition, several ZrO2 crystals reprecipitated at the interface with an elongated shape during cooling (Fig. 1b). The interface between the ceramic partially penetrated by the glass and the glass alone is straight and well de®ned because the overall penetration process is governed by the glass mass transfer and not by the ceramic grains mobility. If the ceramic grains moved from their original position, the interface with the glass would become irregular, as for example in the case of polycrystalline Si3 N4 in contact with a silicate glass [14].
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Fig. 1. (a) EDS element pro®le of sample Z1 (Table 2) across the interface between the residual glass and the ceramic partially penetrated by the glass. (b) Detail of the interface region; the arrows indicate the elongated ZrO2 grains precipitated from the glass at the interface.
The glass penetrated into zirconia in a similar way also during the bonding experiments. The extent of the penetration is function of the starting glass amount, bonding temperature and soaking time, as shown in the plot of Fig. 2a. The microstructural analysis of the interface reveals the presence of a small amount of residual glass only in sample Z5 and Z6 (Table 2) heat treated for 15 min with a starting glass amount equivalent to a thickness of 50 and 30 lm, respectively. These samples, polished only up to 15 lm, have an irregular interfacial microstructure with the ®nal glass thickness
Fig. 2. (a) Plot of the glass penetration in ZrO2 in function of the starting glass interlayer thickness. (b) Residual glass with a thickness of 4±5 lm large ZrO2 crystals of 3±4 lm in sample Z6. (c) Interfacial microstructure of sample Z3.
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varying from about 0 to 4 lm (Fig. 2b). Large, spherical ZrO2 crystals with a diameter up to 10 times greater than the diameter of the grains in the ceramic bulk, formed in the glass interlayer of sample Z6. These grains grew in the asperities and holes of the ZrO2 surfaces, left by the rough polishing, that were not immediately wet by the melted glass. The presence of the so-formed closed voids between the two ceramic surfaces promoted the abnormal grain growth of ZrO2 polycrystals by surface diusion mechanisms. Due to the limited number of these grains, the X-ray diractometry could not determine their crystallographic nature but, considering their dimensions, they are probably cubic since cubic grains are 6±8 times greater than tetragonal grains. Cubic ZrO2 may form in yttria-tetragonal-zirconia-polycrystal ceramics under certain experimental conditions as for example in the above described closed voids at high temperature [7,8]. The ZrO2 grains have probably acquired the spherical shape exhibited in Fig. 2 during the soaking stage at high temperature, when the liquid glass ¯ows through the two ceramic surface and surrounds the grains. In presence of a liquid phase, the grains take the lower energy shape, which is spherical for zirconia. In the vicinity of glass pockets in sintered ZrO2 the grains are in fact more spherical [7,8] as a result of the sphering force that develops in the presence of a liquid phase in many ceramic materials [5,24]. In the other samples (Z1±Z4 of Table 2) no large crystalline ZrO2 particles were observed at the interface because the ceramic surfaces, polished up to 1 lm were wet simultaneously by the glass (Fig. 2c). Under this condition no driving force for abnormal grain growth developed. The consequent interface is characterized by a homogeneous distribution of quasi-spherical zirconia grains, each surrounded by a ``grain boundary'' glassy phase. In this region the grain size of ZrO2 particles is comparable to the grain size of the as-sintered material (0.7±1 lm) but with a more spherical shape resulting from the sphering force. The two ceramic pieces appear so intimately joined (Fig. 2c), that the position of the original interface can be determined only in the vicinity of a residual porosity which can be used as a marker. In these samples neither new phases formed nor the residual glass crystallized. The homogeneous mixture at the interface of the ceramic grains and glass probably inhibited the crystallization during cooling. Al2 O3 ±ZrO2 composite/CaO±Al2 O3 ±SiO2 glass system The behavior of the composite in contact with the glass is dierent compared to ZrO2 . The contact angle during a wetting experiment at 1500°C is similar h < 10° but the glass did not penetrate deeply in the ceramic bulk (Fig. 3a). The ceramic at the interface partially dissolved in the glass, in particular it was depleted of zirconia. As a consequence a new interlayer of about 20 lm grew between the composite and the glass. This layer is formed mainly by alumina and glass. The glass in fact ¯owed through the grain boundaries, promoted the preferential dissolution of zirconia and ®lled the empty voids. On top of this interlayer a row of Al2 O3 grains precipitated from the glass during cooling. These additional grains have an elongated shape and can be distinguished from the grains of the ceramic bulk because they formed on top of the original ceramic surface. The composition of the residual glass changed
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Fig. 3. (a) Wetting experiment, interface of sample AZ. Bonding experiments, (b) sample AZ6, (c) sample AZ2.
compared to the original one. It contains 4±5 wt.% of ZrO2 and the Al2 O3 content increased from 30 to 38 wt.%. The new glass composition is the result of the dissolution of the ceramic in contact with the glass. The composition of the glass is constant along all the thickness, i.e. the glass saturated with ZrO2 and Al2 O3 in 30 min at 1500°C. The ®nal glass composition approximates the eutectic one in the ternary phase diagram of Al2 O3 ±SiO2 ±CaO and this have limited the crystallization. The bonding tests were performed at 1450°C, a temperature higher than the melting point of the glass (1430°C) but lower than 1500°C in order to prevent the Al2 O3 interlayer formation. In all the tested samples (AZ2±AZ7 of Table 2) the glass wets extensively the ceramic and ¯ows along the interface leading to a glassy or partially crystallized interlayer (Fig. 3b and c). The thickness of this interlayer is function of the starting glass amount and of the polishing degree of the ceramic. In the case of surfaces polished with diamond paste up to 1 lm, the ®nal interlayer thickness is equal to the original thickness of the glass, whereas in the samples polished only up to 15 lm the interlayer is slightly thinner because the glass ®lls the super®cial roughness. As expected, only a limited Al2 O3 precipitation occurred at this temperature along the two ceramic interfaces. The extent of the glass interlayer crystallization depended on the starting glass thickness and on the soaking time at high temperature. The in¯uence of the interlayer thickness is highlighted in Fig. 3b and c, where the ®nal microstructure of samples AZ6 and AZ2, characterized by a starting glass thickness of 30 and 50 lm, respectively, is compared. The composition of the large and elongated crystals of sample AZ2 corresponds to the composition of anorthite CaO Al2 O3 2SiO2 , as assessed by EDX analysis and con®rmed by X-ray diractometry performed on the open surfaces of the bars fractured during the bending tests. These elongated grains grew along the c-axis, perpendicularly to the interface. Crystallization of anorthite is in agreement with the original composition of the glass, which falls in the primary phase ®eld of this phase in the ternary diagram SiO2 ±Al2 O3 ±CaO. Zirconia on the other hand dissolved from the ceramic composite and reprecipitated during cooling. Clusters and chains of small spherical ZrO2 grains grew with a dendritic
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shape mainly at the edges of the anorthite grains (indicated with the arrows in Fig. 3c). Zirconia is a nucleating agent for crystallization in glasses [25], and the small crystals which precipitated in the glass promoted the crystallization of anorthite. With short soaking times, a partial crystallization occurred only in samples with a glass interlayer thickness of 30 and 50 lm polished up to 15 lm (AZ6 and AZ7, respectively). In these samples crystallization occurred only in the central part of the interlayer where the residual heat during cooling is higher and the glass has more time to crystallize. In the sample polished up to 1 lm and with a glass interlayer thickness of 20 lm (AZ5), no crystallization occurred. The above reported results evidence that crystallization phenomena inside the glass interlayer depend on the processing parameters and on the quantity of glass used as interlayer. The surface roughness on the contrary did not aect the crystallization. Mechanical properties The mechanical properties of the joints obtained with the two ceramic/glass systems were characterized with two dierent tests. In the ®rst test the adhesion at the ceramic/glass interface was evaluated through the indentation of polished samples with a Vickers tip loaded with 1±5 kg. No crack developed along the ceramic/glass interface. The joints were also characterized with the 3-point bending tests (Table 2). The results showed, in case of both systems, an average strength value higher than 150 MPa, but with large values of standard deviation. In the case of sample Z3, for example, the higher and lower values obtained are 313 and 86 MPa, respectively, and in sample AZ4 these values are 299 and 58 MPa. The large scattering is due to the presence of pores and regions with no glass at the interface. The pores are present in all the tested bars, but are larger in those exhibiting the lower strength values. The pores probably forms at high temperature when the melted glass starts to ¯ow. If closed pores surrounded by the glass form at the interface, the air within them may not escape because the inner pressure is counterbalanced by the capillary force which drives the glass wetting of the ceramic surfaces. In this case the good glass wetting inhibits the pore closure. All the tested samples have shown the presence of residual pores which hiddened the in¯uence of other factors on the mechanical resistance of the joints. For example the eect of the glass thickness, of the surface ®nishing and of the crystallization of the elongated anorthite crystals described in the previous section, can hardly be quanti®ed. The results of the mechanical tests have shown that a glass interlayer can be used to produce joints between oxide ceramics with high ¯exural strength only when the glass wets continuously the ceramic surface. If pores form, the obtained joint is not reliable. Further work is needed to control and eliminate the porosity formation along the interface. One way can be represented by a short pre-bonding cycle at T > Tglass melting of only the ceramic part coated with the glass, which therefore would melt and coat uniformly the ceramic surface. The adhesion process would occur without the pore formation in the subsequent bonding cycle.
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Acknowledgements The authors thank Monica Ferraris and Milena Salvo of the Polytecnic of Turin, Department of Materials Science and Chemical Engineering, for the preparation of the silicate glass and for the useful discussions; Stefano Guicciardi and Cesare Melandri for the mechanical testing. This work was carried out under the aegis of the National Programme PF MSTA II (Progetto Finalizzato Materiali Speciali per Tecnologie Avanzate) of the National Research Council, CNR.
References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11]
[12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22] [23] [24] [25]
Brydson, R., Chen, S.-C., Riley, F. L., Milne, S. J., Pan, X., & R uhle, M. (1998). J Am Ceram Soc 81(2), 369. Flaitz, P. L., & Pask, J. A. (1987). J Am Ceram Soc 70(7), 449. Shaw, T. M., & Duncombe, P. R. (1991). J Am Ceram Soc 74(10), 2495. Kim, J.-J., & Harmer, M. P. (1998). J Am Ceram Soc 81(1), 205. Ramamurthy, S., Mallamaci, M. P., Zimmerman, C. M., Carter, C. B., Duncombe, P. R., & Shaw, T. M. (1996). JMSA 2(3), 113. Kim, Y.-M., Hong, S.-H., & Kim, D.-Y. (2000). J Am Ceram Soc 83(11), 2809. Butler, E. P. (1985). Mat Sci Tech 6(1), 417. R uhle, M., Claussen, N., & Heuer A. H. (1984). Advance in Ceramics (Vol. 12, p. 352). The American Ceramic Society: Ohio. Bellosi, A., Galassi, C., & Guicciardi, S. (1990). J Mat Sci 25, 4331. Choi, H.-J., Kim, G.-H., & Lee, J.-G. (2000). J Am Ceram Soc 83(11), 2821. Nakamura, M., Shigematsu, I., Yamada, Y., Nakanishi, M., Shimojima, K., & Shigegaky, Y. (1998). In K. Niihara (Ed.). Proceedings of the Sixth International Symposium on Ceramic Materials and Components for Engines (p. 305). Kanzaki, S., & Tabata, H. (1983). J Jpn Ceram Soc 91(11), 520. Bellosi, A., & Esposito, L. (1995). In B. S. Tranchina & A. Bellosi (Eds.). Proceedings of the Fourth Euroceramics (Vol. 9, p. 205). Faenza: Italy. Esposito, L., Bellosi, A., & Landi, S. (1999). J Am Ceram Soc 82(12), 3597±3604. Klomp, J. K., & Batden, Th. P. J. (1965). Proc Symp Brit Ceram Res Assoc 48, 129 (Special Publications). Tomsia, A. P., Glaeser, A. M., & Moja, J. S. (1995). Key Eng Mat 111±112, 191. Zdaniewski, W. A., Shah, P. M., & Kirchner, H. P. (1987). Adv Ceram Mater 2(3A), 204. Phase Diagrams for Ceramists, The American Ceramic Society Inc., Fig. 631 (1964). Hasegawa, Y. (1984). Glasstech Ber 57(7), 177. Hasegawa, Y. (1986). Glasstech Ber 59(2), 53. Hasegawa, Y. (1986). Glasstech Ber 59(7), 189. Esposito, L., Bellosi, A., Guicciardi, S., & De Portu, G. (1998). J Mat Sci 33, 1827. Sciti, D., Bellosi, A., & Esposito, L. (2001). J Europ Ceram Soc 21(1), 45. Park, H.-H., Kang, S.-J. L., & Yoon, D. N. (1986). Metall Trans A 17, 325. Beall, G. H., Duke, D. A. (1983). Glass science and technology (Vol. I). In D. R. Uhlmann & N. J. Kreidl (Eds.). Glass Ceramic Techniques (p. 403). Academic Press: New York.
www.elsevier.com/locate/scriptamat
Joining of ceramic oxides by liquid wetting and capillarity Laura Esposito and Alida Bellosi* IRTEC CNR, Research Institute for Ceramics Technology, via Granarolo 64, 48018 Faenza RA, Italy
Received 5 February 2001; accepted 12 June 2001
Abstract A bonding technique based on the use of a glass powder interlayer is studied. Examples are presented for the systems formed by ZrO2 (3 mol%Y2 O3 ) and the composite 80Al2 O3 ±20ZrO2 (wt.%) and a silicate glass as interlayer. The microstructure and mechanical resistance of the joints are also reported. Ó 2001 Published by Elsevier Science Ltd. on behalf of Acta Materialia Inc. Keywords: Bonding; Ceramics; Oxides; Glasses; Interfaces
Introduction Polycrystalline oxide materials like Al2 O3 or ZrO2 and related composites have been widely investigated in the last decades in terms of microstructure [1±6] and thermomechanical properties [7±9]. Most of the characteristics of these oxides are related and governed by the microstructure, i.e. the mean grain size and distribution, the composition and thickness of the intergranular phase, etc. These features play a key role in determining the ®nal properties of the material and depend primarily on the speci®c cations segregation (from impurities and/or sintering aids) occurring during sintering and on the consequent interfacial force balance between the grains [10]. The interactions occurring at high temperature among the phases that come in contact are fundamental also to produce adhesion between two dierent materials. The wetting on ¯owing capability of the intergranular phase, is a key factor for example for the production of bonds through the use of a pressure applied on two ceramic pieces in contact [11±14]. As an alternative to pressure-assisted joining methods, a glass or glass±ceramic interlayer placed between two ceramic pieces can be used to promote adhesion [15±17]. *
Corresponding author. E-mail address: [email protected] (A. Bellosi).
1359-6462/01/$ - see front matter Ó 2001 Published by Elsevier Science Ltd. on behalf of Acta Materialia Inc. PII: S 1 3 5 9 - 6 4 6 2 ( 0 1 ) 0 1 0 9 2 - 2
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No pressure is needed with this technique and, if the glass composition is properly selected, the new phases which may form at the interface are not deleterious for the mechanical resistance of the joint. The optimal glass composition is function of the ceramic material, in particular: ± the ceramic must be chemically and thermomechanically compatible with the glass; ± the glass must wet the ceramic; ± the glass must melt below the temperature at which the ceramic starts to degrade (i.e., below the sintering temperature); ± the glass must have a relatively low viscosity in order to ¯ow easily between the two ceramic parts. In the following, two systems exhibiting a dierent behavior during the bonding cycle are presented. The same Ca±Al±silicate glass is used as interlayer for the composite Al2 O3 ±ZrO2 and for ZrO2 (3 mol%Y2 O3 ). The penetration of the glass in the ceramic and the mobility and solubility of the ceramic grains in the glass drive the ®nal interfacial microstructure. Experimental The compositions and principal characteristics of the ceramic materials are reported in Table 1. The glass used as interlayer (prepared at the Polytechnic of Torino, Italy) has the composition SiO2 39, Al2 O3 30, CaO 31 (wt.%), which falls in the primary phase ®eld for anorthite CaO Al2 O3 2SiO2 [18], has a melting point of 1430°C, as resulted from the heating microscope analysis, a coecient of thermal expansion of 6:6 10 6 K 1 [19±21] and a density of 2.76 g/cm3 . The ceramic surfaces kept in contact during the bonding cycles were previously polished with diamond paste with two dierent grades of surface ®nishing: (i) 1 lm diamond paste was used when the microstructural evolution of the two surfaces in contact and the ®nal glass thickness were analyzed; (ii) a polishing level of 15 lm is adopted to evaluate the feasibility and resistance of joints through an experimental procedure which can be compatible with industrial processes. The glass amount required to obtain the desired thickness was calculated considering the surface area of the ceramic piece and the density of the glass. The glass powder was sieved below 25 lm, weighed, ultrasonically dispersed in ethanol and deposited on the polished and cleaned surface of one ceramic piece. The other ceramic part was placed on top of the deposited glass and the assembled sample was kept at 60°C for 24 h to remove completely the ethanol. Joints were obtained using a laboratory furnace in air, under the experimental conditions reported in Table 2. The interfacial microstructure of Table 1 Compositions, sintering conditions and some characteristics of the materials Material
Sintering cycle (°C h)
Density gr/cm3 (%T.D.)
a (10 6 /K)
r (MPa)
ZrO2 (3 mol% Y2 O3 ) Al2 O3 ±ZrO2 (3 mol% Y2 O3 )
1500 1 1600 1
6.04 (99.8) 4.22 (98.4)
10 8.5
600 20 430 30
a: thermal expansion coecient (20±1200°C), r: 4-pts R.T. ¯exural strength. Experimental procedure described elsewhere for Al2 O3 ±ZrO2 [9], ZrO2 [22,23]. The composition of the composite is 80Al2 O3 ±20ZrO2 (wt.%).
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Table 2 Bonded systems (ZX for ZrO2 =glass joining tests, AZX for Al2 O3 ±ZrO2 =glass joining tests); bonding conditions and 3-pts R.T. ¯exural strength (r) values of the joints. A Anorthite CaO Al2 O3 2SiO2 Sample
Cycle (°C min)
Starting glass thickness (lm)
Polishing level (lm)
Glass penetration (lm)
New crystal phases in glass
Final glass thickness (lm)
r (MPa)
Z1 Z2 Z3 Z4 Z5 Z6
1500 30 1500 30 1500 30 1450 30 1450 15 1450 15
± 20 50 50 50 30
1 1 1 1 15 15
450 200 350 200 200 gl.
ZrO2 None None None None None
± No glass No glass No glass 10 0±3
± ± 173 88 99 56 47 29 ±
AZ1 AZ2 AZ3 AZ4 AZ5 AZ6 AZ7
1500 30 1450 30 1450 30 1450 30 1450 15 1450 15 1450 15
± 50 20 50 20 30 50
1 1 1 15 1 15 15
ZrO2 , Al2 O3 ZrO2 , A None ZrO2 , A None Partiala Partiala
± 50 20 40 10 25±30 25±30
± 148 39 152 67 191 115 156 82 ± ±
a
Crystallization of ZrO2 and anorthite occurred in the central part of sample, not at the edges.
the joints was analyzed by optical microscopy (OM), scanning electron microscopy (SEM) and energy dispersive spectroscopy (EDX). The fracture strength of the bonds was evaluated in 3-point bending on bars 14:0 2:0 1:5 mm3 on a jig of 11 mm span with a crosshead speed of 0.5 mm/min and calculated with the usual bending formula. Results ZrO2 =CaO±Al2 O3 ±SiO2 glass system A wetting experiment (Z1) and various joining cycles (Z2±Z6) were performed with the zirconia/glass system (Table 2). In the wetting experiment the glass wets the ceramic with a contact angle h < 10° and penetrates extensively through the grain boundaries up to a distance from the interface of about 400±500 lm (Fig. 1a). The composition of the residual glass at the interface resulting from the EDX semiquantitative analysis is Al2 O3 30, SiO2 33, CaO 27, ZrO2 8, Y2 O3 2 (wt.%). The content of Ca and Si is less than in the original glass, whereas the content of Zr and Y is higher. These elements diused from the ceramic as a consequence of the partial dissolution of zirconia grains into the glass. In addition, several ZrO2 crystals reprecipitated at the interface with an elongated shape during cooling (Fig. 1b). The interface between the ceramic partially penetrated by the glass and the glass alone is straight and well de®ned because the overall penetration process is governed by the glass mass transfer and not by the ceramic grains mobility. If the ceramic grains moved from their original position, the interface with the glass would become irregular, as for example in the case of polycrystalline Si3 N4 in contact with a silicate glass [14].
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Fig. 1. (a) EDS element pro®le of sample Z1 (Table 2) across the interface between the residual glass and the ceramic partially penetrated by the glass. (b) Detail of the interface region; the arrows indicate the elongated ZrO2 grains precipitated from the glass at the interface.
The glass penetrated into zirconia in a similar way also during the bonding experiments. The extent of the penetration is function of the starting glass amount, bonding temperature and soaking time, as shown in the plot of Fig. 2a. The microstructural analysis of the interface reveals the presence of a small amount of residual glass only in sample Z5 and Z6 (Table 2) heat treated for 15 min with a starting glass amount equivalent to a thickness of 50 and 30 lm, respectively. These samples, polished only up to 15 lm, have an irregular interfacial microstructure with the ®nal glass thickness
Fig. 2. (a) Plot of the glass penetration in ZrO2 in function of the starting glass interlayer thickness. (b) Residual glass with a thickness of 4±5 lm large ZrO2 crystals of 3±4 lm in sample Z6. (c) Interfacial microstructure of sample Z3.
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varying from about 0 to 4 lm (Fig. 2b). Large, spherical ZrO2 crystals with a diameter up to 10 times greater than the diameter of the grains in the ceramic bulk, formed in the glass interlayer of sample Z6. These grains grew in the asperities and holes of the ZrO2 surfaces, left by the rough polishing, that were not immediately wet by the melted glass. The presence of the so-formed closed voids between the two ceramic surfaces promoted the abnormal grain growth of ZrO2 polycrystals by surface diusion mechanisms. Due to the limited number of these grains, the X-ray diractometry could not determine their crystallographic nature but, considering their dimensions, they are probably cubic since cubic grains are 6±8 times greater than tetragonal grains. Cubic ZrO2 may form in yttria-tetragonal-zirconia-polycrystal ceramics under certain experimental conditions as for example in the above described closed voids at high temperature [7,8]. The ZrO2 grains have probably acquired the spherical shape exhibited in Fig. 2 during the soaking stage at high temperature, when the liquid glass ¯ows through the two ceramic surface and surrounds the grains. In presence of a liquid phase, the grains take the lower energy shape, which is spherical for zirconia. In the vicinity of glass pockets in sintered ZrO2 the grains are in fact more spherical [7,8] as a result of the sphering force that develops in the presence of a liquid phase in many ceramic materials [5,24]. In the other samples (Z1±Z4 of Table 2) no large crystalline ZrO2 particles were observed at the interface because the ceramic surfaces, polished up to 1 lm were wet simultaneously by the glass (Fig. 2c). Under this condition no driving force for abnormal grain growth developed. The consequent interface is characterized by a homogeneous distribution of quasi-spherical zirconia grains, each surrounded by a ``grain boundary'' glassy phase. In this region the grain size of ZrO2 particles is comparable to the grain size of the as-sintered material (0.7±1 lm) but with a more spherical shape resulting from the sphering force. The two ceramic pieces appear so intimately joined (Fig. 2c), that the position of the original interface can be determined only in the vicinity of a residual porosity which can be used as a marker. In these samples neither new phases formed nor the residual glass crystallized. The homogeneous mixture at the interface of the ceramic grains and glass probably inhibited the crystallization during cooling. Al2 O3 ±ZrO2 composite/CaO±Al2 O3 ±SiO2 glass system The behavior of the composite in contact with the glass is dierent compared to ZrO2 . The contact angle during a wetting experiment at 1500°C is similar h < 10° but the glass did not penetrate deeply in the ceramic bulk (Fig. 3a). The ceramic at the interface partially dissolved in the glass, in particular it was depleted of zirconia. As a consequence a new interlayer of about 20 lm grew between the composite and the glass. This layer is formed mainly by alumina and glass. The glass in fact ¯owed through the grain boundaries, promoted the preferential dissolution of zirconia and ®lled the empty voids. On top of this interlayer a row of Al2 O3 grains precipitated from the glass during cooling. These additional grains have an elongated shape and can be distinguished from the grains of the ceramic bulk because they formed on top of the original ceramic surface. The composition of the residual glass changed
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Fig. 3. (a) Wetting experiment, interface of sample AZ. Bonding experiments, (b) sample AZ6, (c) sample AZ2.
compared to the original one. It contains 4±5 wt.% of ZrO2 and the Al2 O3 content increased from 30 to 38 wt.%. The new glass composition is the result of the dissolution of the ceramic in contact with the glass. The composition of the glass is constant along all the thickness, i.e. the glass saturated with ZrO2 and Al2 O3 in 30 min at 1500°C. The ®nal glass composition approximates the eutectic one in the ternary phase diagram of Al2 O3 ±SiO2 ±CaO and this have limited the crystallization. The bonding tests were performed at 1450°C, a temperature higher than the melting point of the glass (1430°C) but lower than 1500°C in order to prevent the Al2 O3 interlayer formation. In all the tested samples (AZ2±AZ7 of Table 2) the glass wets extensively the ceramic and ¯ows along the interface leading to a glassy or partially crystallized interlayer (Fig. 3b and c). The thickness of this interlayer is function of the starting glass amount and of the polishing degree of the ceramic. In the case of surfaces polished with diamond paste up to 1 lm, the ®nal interlayer thickness is equal to the original thickness of the glass, whereas in the samples polished only up to 15 lm the interlayer is slightly thinner because the glass ®lls the super®cial roughness. As expected, only a limited Al2 O3 precipitation occurred at this temperature along the two ceramic interfaces. The extent of the glass interlayer crystallization depended on the starting glass thickness and on the soaking time at high temperature. The in¯uence of the interlayer thickness is highlighted in Fig. 3b and c, where the ®nal microstructure of samples AZ6 and AZ2, characterized by a starting glass thickness of 30 and 50 lm, respectively, is compared. The composition of the large and elongated crystals of sample AZ2 corresponds to the composition of anorthite CaO Al2 O3 2SiO2 , as assessed by EDX analysis and con®rmed by X-ray diractometry performed on the open surfaces of the bars fractured during the bending tests. These elongated grains grew along the c-axis, perpendicularly to the interface. Crystallization of anorthite is in agreement with the original composition of the glass, which falls in the primary phase ®eld of this phase in the ternary diagram SiO2 ±Al2 O3 ±CaO. Zirconia on the other hand dissolved from the ceramic composite and reprecipitated during cooling. Clusters and chains of small spherical ZrO2 grains grew with a dendritic
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shape mainly at the edges of the anorthite grains (indicated with the arrows in Fig. 3c). Zirconia is a nucleating agent for crystallization in glasses [25], and the small crystals which precipitated in the glass promoted the crystallization of anorthite. With short soaking times, a partial crystallization occurred only in samples with a glass interlayer thickness of 30 and 50 lm polished up to 15 lm (AZ6 and AZ7, respectively). In these samples crystallization occurred only in the central part of the interlayer where the residual heat during cooling is higher and the glass has more time to crystallize. In the sample polished up to 1 lm and with a glass interlayer thickness of 20 lm (AZ5), no crystallization occurred. The above reported results evidence that crystallization phenomena inside the glass interlayer depend on the processing parameters and on the quantity of glass used as interlayer. The surface roughness on the contrary did not aect the crystallization. Mechanical properties The mechanical properties of the joints obtained with the two ceramic/glass systems were characterized with two dierent tests. In the ®rst test the adhesion at the ceramic/glass interface was evaluated through the indentation of polished samples with a Vickers tip loaded with 1±5 kg. No crack developed along the ceramic/glass interface. The joints were also characterized with the 3-point bending tests (Table 2). The results showed, in case of both systems, an average strength value higher than 150 MPa, but with large values of standard deviation. In the case of sample Z3, for example, the higher and lower values obtained are 313 and 86 MPa, respectively, and in sample AZ4 these values are 299 and 58 MPa. The large scattering is due to the presence of pores and regions with no glass at the interface. The pores are present in all the tested bars, but are larger in those exhibiting the lower strength values. The pores probably forms at high temperature when the melted glass starts to ¯ow. If closed pores surrounded by the glass form at the interface, the air within them may not escape because the inner pressure is counterbalanced by the capillary force which drives the glass wetting of the ceramic surfaces. In this case the good glass wetting inhibits the pore closure. All the tested samples have shown the presence of residual pores which hiddened the in¯uence of other factors on the mechanical resistance of the joints. For example the eect of the glass thickness, of the surface ®nishing and of the crystallization of the elongated anorthite crystals described in the previous section, can hardly be quanti®ed. The results of the mechanical tests have shown that a glass interlayer can be used to produce joints between oxide ceramics with high ¯exural strength only when the glass wets continuously the ceramic surface. If pores form, the obtained joint is not reliable. Further work is needed to control and eliminate the porosity formation along the interface. One way can be represented by a short pre-bonding cycle at T > Tglass melting of only the ceramic part coated with the glass, which therefore would melt and coat uniformly the ceramic surface. The adhesion process would occur without the pore formation in the subsequent bonding cycle.
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Acknowledgements The authors thank Monica Ferraris and Milena Salvo of the Polytecnic of Turin, Department of Materials Science and Chemical Engineering, for the preparation of the silicate glass and for the useful discussions; Stefano Guicciardi and Cesare Melandri for the mechanical testing. This work was carried out under the aegis of the National Programme PF MSTA II (Progetto Finalizzato Materiali Speciali per Tecnologie Avanzate) of the National Research Council, CNR.
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