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The crystallisation of glasses
Journal of Non-Crystalline Solids 52 (1982) 67-76 North-Holland Publishing Company
67
THE CRYSTALLISATION OF GLASSES P.W. M c M I L L A N Department of Physics, University of Warwick, Coventry, England
The importance of studies of glass crystallisation both scientifically and technologically is emphasised and reference is made to structural and kinetic effects that govern crystallization. The application of controlled crystallisation in the formation of glass-ceramics is discussed with reference to the use of nucleating agents. The role of crystal growth in influencing the development of glass-ceramic microstructures and the effects of minor glass constituents and of heat-treatment atmosphere are discussed. In the case of surface-crystallised glasses it is shown that the crystal nucleation density is strongly dependent on the state of the glass surface prior to the crystallisation heat-treatment. Methods of investigating glass crystallisation are discussed with particular reference to electrical measurements and thermal analytical techniques.
1. Introduction The investigation of the crystallisation of glasses is of importance both scientifically and technologically. Glass formation itself implies the avoidance of crystallisation during the cooling of a melt. On the other hand, the production of glass-ceramics requires the achievement of controlled crystallisation leading to the formation of the desirable fine-grained microstructure while avoiding uncontrolled crystallisation during cooling and shaping of the material.
2. Glass formation and devitrification There are two approaches to the problem of glass formation which may be termed the structural and kinetic theories. The former attempts to predict glass formation based u p o n the atomic arrangements thought to exist in the melt at temperatures just above the liquidus temperature. The kinetic theory is based on the consideration of the rate at which the disordered liquid structure can be transformed into a regular crystal lattice as governed by diffusion processes. Clearly, the kinetic factors that govern glass formation must be related to atomic arrangements and bonding within the melt and hence to structure. It is true, however, that present knowledge of the structure of liquids and in particular of the complex liquids often utilised for glass formation is fragmentary and considerable research is needed into this field. Thus structural theories are essentially qualitative while the kinetic theory is more quantitative. A structural theory that has gained wide acceptance is the r a n d o m network 0 0 2 2 - 3 0 9 3 / 8 2 / 0 0 0 0 - 0 0 0 0 / $ 0 2 . 7 5 © 1982 N o r t h - H o l l a n d
P. W. McMillan / The crystallisation of glasses
68
theory proposed by Zachariasen [1]. In considering the limits of glass formation in various systems some possible limitations of the theory become apparent. The structure of vitreous silica is envisaged as a continuous non-periodic network of SiO4 tetrahedra linked together at all four corners by bridging oxygens. The introduction of an alkali metal oxide to form a binary silicate composition will introduce pairs of non-bridging oxygens which are linked to only one silicon and in the limit, at the orthosilicate composition (2R20 • SiO 2) all of the oxygens will become non-bridging and a continuous silicate network will no longer be possible. It is likely, however, that glass formation would become increasingly improbable when the proportion of the second oxide exceeds that given by the metasilicate composition (R20-SiO2). At this composition, the average number of bridging oxygens per tetrahedron is two. Therefore, rings or long silicate chains could still occur and glass formation would perhaps still be possible. Further increase of the content of the added oxide would cause decrease of average chain length and disruption of rings bringing about a progressive fall in glass viscosity. This would favour enhanced crystal nucleation and growth rates and increase the probability of devitrification. In the alkali oxide-silica systems it is found that glass formation does not readily occur for compositions where the average number of bridging oxygens per SiO4 group is less than two. The glass forming limits for the L i 2 0 - , N a 2 0 and K 2 0 - S i O 2 systems are 40, 47 and 50 mol% respectively. These findings also underline the fact that one cannot simply consider the bonding within the network-forming component (silica in this case). Clearly, the added modifier cation plays an important role and the more restricted range of glass formation in the L i 2 0 - S i O 2 system is most likely related to the higher ionic field strength of this ion enabling it to exert a greater ordering effect on surrounding oxygen ions in the melt. The important role of modifier cations is even more apparent when the so-called "invert" glasses are considered. Moore and mcMillan [2] showed that glass formation in the L i 2 0 - M g O - S i O 2 system was possible for a silica content as low as 35 mol% and Trap and Stevels [3] observed glass formation for SiO 2 contents as low as 40 tool% in complex silicate glasses. Assuming a homogeneous distribution of non-bridging oxygens, the melt structure in these cases would consist of isolated pairs of SiO4 tetrahedra. To explain glass formation in these cases, it is necessary to postulate structural roles for the Mg z+ and other alkaline earth ions in which they serve to link together SiO 4 groups in a manner not envisaged by the random network theory. Similarly, in the case of PbO-SiO z melts in which glass formation occurs up to the 2PbO. SiO 2 composition, and thus where isolated SiO 4 tetrahedra would be present, a structural role in which Pb 2+ ions form bridges between adjacent tetrahedra appears to be necessary. The basis of kinetic theories of glass formation is the Johnson-Male-Avrami equation: X t
=
1 - exp( - Kt ~),
P.W. McMillan / The crvstallisation of glasses
69
where x, is the volume fraction crystallised after time t and n is the Avrami exponent and is equal to 4 for constant nucleation and growth rates. The rate constant K may be written as: ('n'/3)U3I where U is the crystal growth rate and 1 is the nucleation rate. Thus the failure of a melt to crystallise during cooling may imply that either the nucleation rate or the crystal growth rate or both are too low. Expressions for the nucleation rate I and the growth rate U can be approximated as follows:
N W*/RT) e x p ( - AG'/RT), e x p ( - AG"/RT)[1 - exp(AG/RT)],
I = constant X e x p ( U = constant x
where W* is termed the thermodynamic barrier to nucleation after Turnbull and Cohen [4] and AG is the bulk free energy of crystallisation. It will be noted that both expressions contain kinetic terms which are functions of an activation energy: AG', in the case of nucleation and ,AG" in the case of growth. These are both related to diffusion processes but in general they are not equal. In the case of growth, diffusion of atoms over long distances to the growth interface may be required but with regard to nucleation, atomic movements close to the nucleus centre are involved. It is likely that the activation energies AG' and AG" are related to that for viscous flow and therefore that rates of nucleation and growth will be influenced by the coefficient of viscosity ~ and its dependence on temperature. In fact a simple expression for growth rate U, that is found to hold with reasonable accuracy for a number of glass forming systems at temperatures not too far below the liquidus temperature is: U = constant x
AT/q,
where AT is the undercooling below the liquidus temperature. Thus critical factors in glass formation will be the occurrence of a high viscosity at the liquidus and a rapid rate of increase of this parameter as the temperature is further reduced. Clearly the viscosity of the melt will be dependent upon structural effects. Increase of the concentration of non-bridging oxygens leading to disruption of the glass-forming network will cause reduction of viscosity. The role of some modifying cations, however, may be to counterbalance this effect to some extent because of ionic bonding to non-bridging oxygens.
3. Controlled crystallisation The application of controlled crystallisation in the production of glassceramics requires that attention be paid to both nucleation and growth rates. If the former rate is too low, growth will take place from too few centres and a coarse-grained microstructure is likely to result. Also if the crystal growth rate is too high, coarsening of the microstructure can result.
70
P. IV.. McMillan / The crystallisation of glasses
Crystal nucleation may be either homogeneous or heterogeneous but in glasses it is found that homogeneous nucleation is comparatively rare except for some glasses of simple composition. In most cases, heterogeneities must be present since these can lower the activation energy of nucleation if wetting of the heterogeneity by the precipitated phase occurs. In the devitrification of conventional glasses, crystallisation is often nucleated by heterogeneities present on the surface of the glass, such as dust particles or in some cases scratches. While surface nucleation can be utilised for the production of high strength surface crystallised glasses (discussed later) the production of glassceramics requires the generation of a high density of nucleation sites within the bulk of the material. For this purpose, nucleating agents are incorporated into the parent glass compositions. These include metals such as copper, silver and gold and the platinum group metals. In these cases, the metals are precipitated as colloidal dispersions by controlled heat-treatment and these heterogeneously nucleate crystalline phases such as lithium silicates. Undoubtedly in these cases, efficient nucleation is favoured by the occurrence of epitaxial growth of the silicate crystal phase upon the metallic colloid. The most important group of nucleating agents comprises certain oxides including TiO 2, P205 and ZrO 2. The mode of operation of these is more problematical than for the metallic nucleating agents. It is found for many compositions utilising these nucleating agents that the glass separates into two vitreous phases prior to the occurrence of crystal nucleation. It is of interest therefore to consider to what extent this favourably influences crystal nucleation. In the case of TiO 2 the work of Maurer [5] indicated that as a first step a titania-rich phase separates out. At a later stage in the heat-treatment, crystalline titanates are formed possibly by homogeneous nucleation of the TiO2-rich phase. The titanate phase then serves to nucleate silicate and aluminosilicate phases. A somewhat different picture emerges in the case of P205. This oxide undoubtedly enhances prior glass separation as shown by James and McMillan [6], Tomozawa [7] and Matusita et al. [8]. However, experimental evidence does not support the idea that the phase-separated particles serve to provide heterogeneous nucleation sites. McMillan [9] found no simple relationship between the number density of crystal nuclei and the interfacial area or number density of the phase separated particles. Since the occurrence of prior phase separation undoubtedly resulted in a finer-grained ultimate microstructure, it was proposed that the effect of the phase separation was to inhibit the disappearance of nuclei (arising in the matrix phase) by a coarsening mechanism. A further possible effect of P205 is that it inhibits the growth rate of lithium silicate crystals. This was shown by direct measurements of the growth rate of lithium disilicate by McMillan and Phillips [10]. These investigators [11] demonstrated the same effect by measurements of electrical resistivity during the crystallisation process. The reduction of crystal growth rate, brought about by the inclusion of P205, is undoubtedly beneficial in inhibiting coarsening of
P.W. McMillan / The crystallisation of glasses
71
the glass-ceramic microstructure. Small concentrations of other oxides in the glass can also influence crystal growth rates and morphology. Sometimes these inclusions are unintentional, as in the case of water which is almost invariably present in oxide glasses in the form of hydroxyl ions. Hibberd and McMillan [12] showed that the presence of small traces of hydroxyl ions exerted a marked influence on the growth of lithium disilicate. Increasing the OH content from - 0 . 0 1 5 wt% to - 0 . 0 3 wt% caused a displacement of the temperature of maximum crystal growth from 890°C to 870°C. This is due to the effect of hydroxyl ions in reducing the viscosity of silicate glasses [13]. In some cases, the crystallisation kinetics can be affected by the surrounding atmosphere during heat-treatment. Although glass-ceramics would normally be heat-treated in air, there are cases where a non-oxidising or reducing atmosphere may be necessary, as for example when a glass-ceramic is being bonded to an oxidisable metal. If the glass-ceramic contains oxides that can exhibit non-stoichiometry, both nucleation and growth rates may be influenced. The efficiency of TiO 2 as a nucleating agent is severely impaired if heat-treatment is carried out in a nonoxidising atmosphere and in certain glass-ceramics, crystal growth rates can also be affected. Fig. 1 shows results obtained for the growth of a zinc aluminate spinel and it is evident that the heat-treatment atmosphere exerts a marked effect. It is well known that zinc oxide can fairly readily become oxygen-deficient and this evidently inhibits growth of the zinc-containing crystal.
2-0-
A 1-5. E
1-0.
~J 0.5"
900
1000
1100 Temperature ('C.)
1200
1300
Fig. 1. The effects of the heat-treatment atmosphere on the growth of zinc aluminate spinel from a Z n O - A l 2 0 3 - S i O 2 glass: A, oxygen; B, air; C, argon; D, 10% H2/90% N 2.
P. IV.. McMillan / The crystallisation of glasses
72
Table I Relationship between prior surface treatment and surface crystal nucleation density for a ZnO-A1203-SiO 2 glass, heat-treated at 840°C for 10 min. Surface treatment
Nucleation density (/t m - 2 )
220 grit grinding Partial polish Full polish
3.2× l0 2 3.5 x l0 3 l × l0 -3
4. Surface erystailisation of glasses A s a n a l t e r n a t i v e t o b u l k c r y s t a l l i s a t i o n , it h a s b e e n d i s c o v e r e d t h a t c e r t a i n glasses can be heat-treated
to produce
a surface-crystallised
layer resulting in
f
Fig. 2. The effects of prior surface finish on the surface crystal nucleation density of a surfacecrystallisable glass: (a) 200 grit surface finish, (b) polished - 6 p~m diamond finish. Magnification × 600.
P.W. McMillan / The crystallisation of glasses"
73
Table 2 Effects of prior surface treatment upon the strength of a surface-crystallised glass Surface treatment
Mean rupture modulus (MN m 2)
Polished (6 ~ m diamond paste) Polished and free abrasion Polished and 800 grit grinding Special diamond wheel finish
314 388 443 510
considerable improvement in mechanical strength combined in some cases with optical transparency [14]. There are two mechanisms by which improved mechanical strength can arise. One of these involves pre-stressing of the surface layer in compression owing to a thermal expansion mismatch between this layer and the uncrystallised interior of the glass. Secondly, and perhaps more importantly [15] the surface microstructure that develops serves to limit the severity of surface flaws thereby resulting in enhancement of mechanical strength. The achievement of high strength requires the development of a fine-grained surface microstructure and this in turn poses the need for a high density of surface nucleation sites. Investigations have shown that flaws present in the glass surface prior to heat treatment act as nucleation sites. A high density of flaws therefore promotes the formation of a fine-grained microstructure in the surface layer. Adams and McMillan [16] obtained clear evidence that the mechanical strength of a surface crystallised glass could be enhanced by suitable treatment of the glass surface prior to the crystallisation treatment. Table 1 summarises surface nucleation densities determined for a ZnO-A1203-SiO 2 glass. Clearly, the nucleation density is greater for the more severely damaged surface (200 grit grinding). The marked differences in nucleation density brought about by different surface finishes prior to heat treatment are shown in fig. 2. As shown by the results in table 2, considerable enhancement of the strength of surface-crystallised glasses is possible depending on the surface pretreatment. This strength enhancement is related to the development of finer textured surface microstructure resulting from improved surface nucleation density.
5. Methods of investigating crystallisation of glasses A number of techniques are known for the investigation of glass crystallisation and for the study of the development of glass-ceramic microstructures. These include X-ray diffraction used for the identification of crystal phases
P. W. McMillan / The crystallisation of glasses
74
and for the estimation of volume fractions. Optical and electron microscopy, especially the latter, combined with stereological techniques are used to derive useful microstructural parameters such as mean particle sizes, volume fractions of phases and mean interparticle spacings. It has proved possible to relate a number of physical properties of glass-ceramics to microstructural parameters. The application of these techniques has been reviewed elsewhere [17] and will not therefore be discussed further in the present paper. The use of electrical measurements to monitor the progress of crystallisation has already been mentioned and the value of this technique may be further illustrated by the work of Atkinson and McMillan [18]. These investigators studied the crystallisation of a glass essentially of the L i 2 0 - S i O 2 type which was modified by the incorporation of small additions of K20, AI203 and P205, the latter being included as a nucleating agent. Specimens of the glass were heat-treated at temperatures in the range 650 to 900°C for a period of 1 h. Measurements of electrical resistivity made on the specimens after heat-treatment showed that this achieved a maximum after heat-treatment at 750°C. Similarly, the dielectric loss of the material was a minimum after heat-treatment at the same temperature. Data for the power factor, Ktan3, given in fig. 3 illustrate the effect. It was also observed that the mechanical strength of the material was a maximum after heat-treatment at 750°C. These effects are attributed to maximisation of the volume fraction of the crystalline phases and electron microscopy gave support to this finding. The maximum in electrical resistivity and minimum in dielectric losses arising from heat-treatment at
20.
UP
o 10 "5
.'r
t
600
t
[
7OO 8O0 9O0 Heat treatment temperature (°C.)
4
8 12 16 "lime(hours)
20
24
Fig. 3. The effects of heat-treatment temperature on the power factor of a glass-ceramic; measurements made at 300°C and a frequency of 100 kHz. Fig. 4. Heat evolution during isothermal heat-treatment at 730°C for MgO-AI2Oa-B203 glassceramic. A, nucleated at 6 7 0 ° C / 4 h; B, nucleated at 670°C/12 h.
P.W. McMillan / The crystallisation of glasses
75
750°C result from the incorporation of lithium ions into crystalline lithium disilicate in which their mobility is greatly reduced as compared with that in the glass phase. These results and those from similar investigations underline the value of electrical measurements as a sensitive tool for the study of microstructural changes brought about by glass crystallisation. Another valuable class of investigative techniques is represented by thermal analytical methods. These include differential thermal analysis (DTA) and differential scanning calorimetry (DSC). The use of these techniques can be illustrated by the results of Lau and McMillan [19] who investigated the crystallisation of a glass of the MgO-A1203-B203 type. Preliminary trials indicated that conversion of this material to a fine-grained glass-ceramic could be achieved by a nucleation treatment at 670°C and a crystallisation treatment at 730°C. Using a differential scanning calorimeter, the rate of heat evolution at 730°C was determined for specimens that had received various nucleation treatments. Typical results are shown in fig. 4. A striking feature of the results is the change in slope of the curves after heat-treatment at 730°C for about 4 h. It was also noted that the glass-ceramics produced were transparent for heat-treatments up to 4 h but became opaque thereafter. The change in slope of the DSC curves was taken to indicate that two crystalline phases involving
2"
i/
..2"
~0 i5 i-2-
-4" A, / S 600
700
800
Temperature ('C)
900
9 10 Int.(sec.)
1
Fig. 5. DTA curves for MgO-AI203-B203 glass ceramic. A, nucleated at 670°C/12 h; B, nucleated at 670°C/4 h; C, no nucleation treatment. Fig. 6. DSC data for MgO-AI203-B203 glass-ceramic. A, nucleated at 670°C/4 h; B, nucleated at 670°C/12 h.
P. W. McMillan / The crystallisation of glasses
76
different kinetics were being formed. DTA studies, illustrated in fig. 5, confirmed that for specimens that had received a nucleation treatment, two crystalline phases were being formed as shown by the occurrence of two exothermic peaks. It was deduced that the lower temperature peak related to the crystal phase responsible for the initial portion of the DSC curve. Further DTA studies showed that the higher temperature peak was shifted to lower temperatures by the use of finer glass powder whereas the lower peak was not affected. This result indicated that the higher temperature phase was surfacenucleated while the other phase was bulk-nucleated. Further analysis of the DSC data confirmed this finding. From plots of the type illustrated in fig. 6 it is possible to derive the Avrami exponent, n, from the slopes of the two straight line regions. It was found that for the lower temperature phase the value of n was 3 + 0.2. This implies a constant growth rate and a nucleation rate of zero at the treatment temperature of 730°C. For the higher temperature phase, the value of n was found to lie between 0.6 and 1, which could imply either the thickening of plates when complete edge impingement has occurred or the growth of needle-shaped crystals. Either possibility is consistent with surface-nucleated crystallisation.
References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13] [14] [15] [16] [17] [18] [19]
W.H. Zachariasen, J. Am. Ceram. Soc. 54 (1932) 3841. H. Moore and P.W. McMillan, J. Soc. Glass Technol. 40 (1956) 66. H.L. Trap and J.M. Stevels, Glastech. Ber. 32K (1959) 51. D. Turnbull and M.H. Cohen, J. Chem. Phys. 29 (1958) 1049. R.D. Maurer, J. Appl. Phys. 33 (1962) 2132. P.F. James and P.W. McMillan, Phil. Mag. 18 (1968). M. Tomozawa, in: Advances in Nucleation and Crystallization in Glasses (American Ceramic Society, 1971) p. 41. K. Matusita, S. Sakka, T. Muki and M. Tashiro, J. Mat. Sci. 10 (1975) 94. P.W. McMillan, Proc. Xth Int. Cong. on Glass 14 (1974) 1. P.W. McMillan and S.V. Phillips, in: Glass-Ceramics (Academic Press, London, 1979) p. 53. S.V. Phillips and P.W. McMillan, Glass Technol. 7 (1965) 121. B. Hibberd and P.W. McMillan, in: Glass-Ceramics (Academic Press, London, 1979) p. 54. A. Chlebik and P.W. McMillan, J. Non-Crystalline Solids 38 (1980) 509. J.S. Olcott and S.D. Stookey, in: Advances in Glass Technology (Plenum, New York, 1962) p. 400. G. Partridge and P.W. McMillan, Glass Technol. 15 (1974) 127. R. Adams and P.W. McMillan, J. Mat. Sci. 17 (1982) 2727. P.W. McMillan, Glass-ceramics (Academic Press, London, 1979). D.I.H. Atkinson and P.W. McMillan, J. Mat. Sci. 11 (1976) 994. J. Lau and P.W. McMillan, to be published.
67
THE CRYSTALLISATION OF GLASSES P.W. M c M I L L A N Department of Physics, University of Warwick, Coventry, England
The importance of studies of glass crystallisation both scientifically and technologically is emphasised and reference is made to structural and kinetic effects that govern crystallization. The application of controlled crystallisation in the formation of glass-ceramics is discussed with reference to the use of nucleating agents. The role of crystal growth in influencing the development of glass-ceramic microstructures and the effects of minor glass constituents and of heat-treatment atmosphere are discussed. In the case of surface-crystallised glasses it is shown that the crystal nucleation density is strongly dependent on the state of the glass surface prior to the crystallisation heat-treatment. Methods of investigating glass crystallisation are discussed with particular reference to electrical measurements and thermal analytical techniques.
1. Introduction The investigation of the crystallisation of glasses is of importance both scientifically and technologically. Glass formation itself implies the avoidance of crystallisation during the cooling of a melt. On the other hand, the production of glass-ceramics requires the achievement of controlled crystallisation leading to the formation of the desirable fine-grained microstructure while avoiding uncontrolled crystallisation during cooling and shaping of the material.
2. Glass formation and devitrification There are two approaches to the problem of glass formation which may be termed the structural and kinetic theories. The former attempts to predict glass formation based u p o n the atomic arrangements thought to exist in the melt at temperatures just above the liquidus temperature. The kinetic theory is based on the consideration of the rate at which the disordered liquid structure can be transformed into a regular crystal lattice as governed by diffusion processes. Clearly, the kinetic factors that govern glass formation must be related to atomic arrangements and bonding within the melt and hence to structure. It is true, however, that present knowledge of the structure of liquids and in particular of the complex liquids often utilised for glass formation is fragmentary and considerable research is needed into this field. Thus structural theories are essentially qualitative while the kinetic theory is more quantitative. A structural theory that has gained wide acceptance is the r a n d o m network 0 0 2 2 - 3 0 9 3 / 8 2 / 0 0 0 0 - 0 0 0 0 / $ 0 2 . 7 5 © 1982 N o r t h - H o l l a n d
P. W. McMillan / The crystallisation of glasses
68
theory proposed by Zachariasen [1]. In considering the limits of glass formation in various systems some possible limitations of the theory become apparent. The structure of vitreous silica is envisaged as a continuous non-periodic network of SiO4 tetrahedra linked together at all four corners by bridging oxygens. The introduction of an alkali metal oxide to form a binary silicate composition will introduce pairs of non-bridging oxygens which are linked to only one silicon and in the limit, at the orthosilicate composition (2R20 • SiO 2) all of the oxygens will become non-bridging and a continuous silicate network will no longer be possible. It is likely, however, that glass formation would become increasingly improbable when the proportion of the second oxide exceeds that given by the metasilicate composition (R20-SiO2). At this composition, the average number of bridging oxygens per tetrahedron is two. Therefore, rings or long silicate chains could still occur and glass formation would perhaps still be possible. Further increase of the content of the added oxide would cause decrease of average chain length and disruption of rings bringing about a progressive fall in glass viscosity. This would favour enhanced crystal nucleation and growth rates and increase the probability of devitrification. In the alkali oxide-silica systems it is found that glass formation does not readily occur for compositions where the average number of bridging oxygens per SiO4 group is less than two. The glass forming limits for the L i 2 0 - , N a 2 0 and K 2 0 - S i O 2 systems are 40, 47 and 50 mol% respectively. These findings also underline the fact that one cannot simply consider the bonding within the network-forming component (silica in this case). Clearly, the added modifier cation plays an important role and the more restricted range of glass formation in the L i 2 0 - S i O 2 system is most likely related to the higher ionic field strength of this ion enabling it to exert a greater ordering effect on surrounding oxygen ions in the melt. The important role of modifier cations is even more apparent when the so-called "invert" glasses are considered. Moore and mcMillan [2] showed that glass formation in the L i 2 0 - M g O - S i O 2 system was possible for a silica content as low as 35 mol% and Trap and Stevels [3] observed glass formation for SiO 2 contents as low as 40 tool% in complex silicate glasses. Assuming a homogeneous distribution of non-bridging oxygens, the melt structure in these cases would consist of isolated pairs of SiO4 tetrahedra. To explain glass formation in these cases, it is necessary to postulate structural roles for the Mg z+ and other alkaline earth ions in which they serve to link together SiO 4 groups in a manner not envisaged by the random network theory. Similarly, in the case of PbO-SiO z melts in which glass formation occurs up to the 2PbO. SiO 2 composition, and thus where isolated SiO 4 tetrahedra would be present, a structural role in which Pb 2+ ions form bridges between adjacent tetrahedra appears to be necessary. The basis of kinetic theories of glass formation is the Johnson-Male-Avrami equation: X t
=
1 - exp( - Kt ~),
P.W. McMillan / The crvstallisation of glasses
69
where x, is the volume fraction crystallised after time t and n is the Avrami exponent and is equal to 4 for constant nucleation and growth rates. The rate constant K may be written as: ('n'/3)U3I where U is the crystal growth rate and 1 is the nucleation rate. Thus the failure of a melt to crystallise during cooling may imply that either the nucleation rate or the crystal growth rate or both are too low. Expressions for the nucleation rate I and the growth rate U can be approximated as follows:
N W*/RT) e x p ( - AG'/RT), e x p ( - AG"/RT)[1 - exp(AG/RT)],
I = constant X e x p ( U = constant x
where W* is termed the thermodynamic barrier to nucleation after Turnbull and Cohen [4] and AG is the bulk free energy of crystallisation. It will be noted that both expressions contain kinetic terms which are functions of an activation energy: AG', in the case of nucleation and ,AG" in the case of growth. These are both related to diffusion processes but in general they are not equal. In the case of growth, diffusion of atoms over long distances to the growth interface may be required but with regard to nucleation, atomic movements close to the nucleus centre are involved. It is likely that the activation energies AG' and AG" are related to that for viscous flow and therefore that rates of nucleation and growth will be influenced by the coefficient of viscosity ~ and its dependence on temperature. In fact a simple expression for growth rate U, that is found to hold with reasonable accuracy for a number of glass forming systems at temperatures not too far below the liquidus temperature is: U = constant x
AT/q,
where AT is the undercooling below the liquidus temperature. Thus critical factors in glass formation will be the occurrence of a high viscosity at the liquidus and a rapid rate of increase of this parameter as the temperature is further reduced. Clearly the viscosity of the melt will be dependent upon structural effects. Increase of the concentration of non-bridging oxygens leading to disruption of the glass-forming network will cause reduction of viscosity. The role of some modifying cations, however, may be to counterbalance this effect to some extent because of ionic bonding to non-bridging oxygens.
3. Controlled crystallisation The application of controlled crystallisation in the production of glassceramics requires that attention be paid to both nucleation and growth rates. If the former rate is too low, growth will take place from too few centres and a coarse-grained microstructure is likely to result. Also if the crystal growth rate is too high, coarsening of the microstructure can result.
70
P. IV.. McMillan / The crystallisation of glasses
Crystal nucleation may be either homogeneous or heterogeneous but in glasses it is found that homogeneous nucleation is comparatively rare except for some glasses of simple composition. In most cases, heterogeneities must be present since these can lower the activation energy of nucleation if wetting of the heterogeneity by the precipitated phase occurs. In the devitrification of conventional glasses, crystallisation is often nucleated by heterogeneities present on the surface of the glass, such as dust particles or in some cases scratches. While surface nucleation can be utilised for the production of high strength surface crystallised glasses (discussed later) the production of glassceramics requires the generation of a high density of nucleation sites within the bulk of the material. For this purpose, nucleating agents are incorporated into the parent glass compositions. These include metals such as copper, silver and gold and the platinum group metals. In these cases, the metals are precipitated as colloidal dispersions by controlled heat-treatment and these heterogeneously nucleate crystalline phases such as lithium silicates. Undoubtedly in these cases, efficient nucleation is favoured by the occurrence of epitaxial growth of the silicate crystal phase upon the metallic colloid. The most important group of nucleating agents comprises certain oxides including TiO 2, P205 and ZrO 2. The mode of operation of these is more problematical than for the metallic nucleating agents. It is found for many compositions utilising these nucleating agents that the glass separates into two vitreous phases prior to the occurrence of crystal nucleation. It is of interest therefore to consider to what extent this favourably influences crystal nucleation. In the case of TiO 2 the work of Maurer [5] indicated that as a first step a titania-rich phase separates out. At a later stage in the heat-treatment, crystalline titanates are formed possibly by homogeneous nucleation of the TiO2-rich phase. The titanate phase then serves to nucleate silicate and aluminosilicate phases. A somewhat different picture emerges in the case of P205. This oxide undoubtedly enhances prior glass separation as shown by James and McMillan [6], Tomozawa [7] and Matusita et al. [8]. However, experimental evidence does not support the idea that the phase-separated particles serve to provide heterogeneous nucleation sites. McMillan [9] found no simple relationship between the number density of crystal nuclei and the interfacial area or number density of the phase separated particles. Since the occurrence of prior phase separation undoubtedly resulted in a finer-grained ultimate microstructure, it was proposed that the effect of the phase separation was to inhibit the disappearance of nuclei (arising in the matrix phase) by a coarsening mechanism. A further possible effect of P205 is that it inhibits the growth rate of lithium silicate crystals. This was shown by direct measurements of the growth rate of lithium disilicate by McMillan and Phillips [10]. These investigators [11] demonstrated the same effect by measurements of electrical resistivity during the crystallisation process. The reduction of crystal growth rate, brought about by the inclusion of P205, is undoubtedly beneficial in inhibiting coarsening of
P.W. McMillan / The crystallisation of glasses
71
the glass-ceramic microstructure. Small concentrations of other oxides in the glass can also influence crystal growth rates and morphology. Sometimes these inclusions are unintentional, as in the case of water which is almost invariably present in oxide glasses in the form of hydroxyl ions. Hibberd and McMillan [12] showed that the presence of small traces of hydroxyl ions exerted a marked influence on the growth of lithium disilicate. Increasing the OH content from - 0 . 0 1 5 wt% to - 0 . 0 3 wt% caused a displacement of the temperature of maximum crystal growth from 890°C to 870°C. This is due to the effect of hydroxyl ions in reducing the viscosity of silicate glasses [13]. In some cases, the crystallisation kinetics can be affected by the surrounding atmosphere during heat-treatment. Although glass-ceramics would normally be heat-treated in air, there are cases where a non-oxidising or reducing atmosphere may be necessary, as for example when a glass-ceramic is being bonded to an oxidisable metal. If the glass-ceramic contains oxides that can exhibit non-stoichiometry, both nucleation and growth rates may be influenced. The efficiency of TiO 2 as a nucleating agent is severely impaired if heat-treatment is carried out in a nonoxidising atmosphere and in certain glass-ceramics, crystal growth rates can also be affected. Fig. 1 shows results obtained for the growth of a zinc aluminate spinel and it is evident that the heat-treatment atmosphere exerts a marked effect. It is well known that zinc oxide can fairly readily become oxygen-deficient and this evidently inhibits growth of the zinc-containing crystal.
2-0-
A 1-5. E
1-0.
~J 0.5"
900
1000
1100 Temperature ('C.)
1200
1300
Fig. 1. The effects of the heat-treatment atmosphere on the growth of zinc aluminate spinel from a Z n O - A l 2 0 3 - S i O 2 glass: A, oxygen; B, air; C, argon; D, 10% H2/90% N 2.
P. IV.. McMillan / The crystallisation of glasses
72
Table I Relationship between prior surface treatment and surface crystal nucleation density for a ZnO-A1203-SiO 2 glass, heat-treated at 840°C for 10 min. Surface treatment
Nucleation density (/t m - 2 )
220 grit grinding Partial polish Full polish
3.2× l0 2 3.5 x l0 3 l × l0 -3
4. Surface erystailisation of glasses A s a n a l t e r n a t i v e t o b u l k c r y s t a l l i s a t i o n , it h a s b e e n d i s c o v e r e d t h a t c e r t a i n glasses can be heat-treated
to produce
a surface-crystallised
layer resulting in
f
Fig. 2. The effects of prior surface finish on the surface crystal nucleation density of a surfacecrystallisable glass: (a) 200 grit surface finish, (b) polished - 6 p~m diamond finish. Magnification × 600.
P.W. McMillan / The crystallisation of glasses"
73
Table 2 Effects of prior surface treatment upon the strength of a surface-crystallised glass Surface treatment
Mean rupture modulus (MN m 2)
Polished (6 ~ m diamond paste) Polished and free abrasion Polished and 800 grit grinding Special diamond wheel finish
314 388 443 510
considerable improvement in mechanical strength combined in some cases with optical transparency [14]. There are two mechanisms by which improved mechanical strength can arise. One of these involves pre-stressing of the surface layer in compression owing to a thermal expansion mismatch between this layer and the uncrystallised interior of the glass. Secondly, and perhaps more importantly [15] the surface microstructure that develops serves to limit the severity of surface flaws thereby resulting in enhancement of mechanical strength. The achievement of high strength requires the development of a fine-grained surface microstructure and this in turn poses the need for a high density of surface nucleation sites. Investigations have shown that flaws present in the glass surface prior to heat treatment act as nucleation sites. A high density of flaws therefore promotes the formation of a fine-grained microstructure in the surface layer. Adams and McMillan [16] obtained clear evidence that the mechanical strength of a surface crystallised glass could be enhanced by suitable treatment of the glass surface prior to the crystallisation treatment. Table 1 summarises surface nucleation densities determined for a ZnO-A1203-SiO 2 glass. Clearly, the nucleation density is greater for the more severely damaged surface (200 grit grinding). The marked differences in nucleation density brought about by different surface finishes prior to heat treatment are shown in fig. 2. As shown by the results in table 2, considerable enhancement of the strength of surface-crystallised glasses is possible depending on the surface pretreatment. This strength enhancement is related to the development of finer textured surface microstructure resulting from improved surface nucleation density.
5. Methods of investigating crystallisation of glasses A number of techniques are known for the investigation of glass crystallisation and for the study of the development of glass-ceramic microstructures. These include X-ray diffraction used for the identification of crystal phases
P. W. McMillan / The crystallisation of glasses
74
and for the estimation of volume fractions. Optical and electron microscopy, especially the latter, combined with stereological techniques are used to derive useful microstructural parameters such as mean particle sizes, volume fractions of phases and mean interparticle spacings. It has proved possible to relate a number of physical properties of glass-ceramics to microstructural parameters. The application of these techniques has been reviewed elsewhere [17] and will not therefore be discussed further in the present paper. The use of electrical measurements to monitor the progress of crystallisation has already been mentioned and the value of this technique may be further illustrated by the work of Atkinson and McMillan [18]. These investigators studied the crystallisation of a glass essentially of the L i 2 0 - S i O 2 type which was modified by the incorporation of small additions of K20, AI203 and P205, the latter being included as a nucleating agent. Specimens of the glass were heat-treated at temperatures in the range 650 to 900°C for a period of 1 h. Measurements of electrical resistivity made on the specimens after heat-treatment showed that this achieved a maximum after heat-treatment at 750°C. Similarly, the dielectric loss of the material was a minimum after heat-treatment at the same temperature. Data for the power factor, Ktan3, given in fig. 3 illustrate the effect. It was also observed that the mechanical strength of the material was a maximum after heat-treatment at 750°C. These effects are attributed to maximisation of the volume fraction of the crystalline phases and electron microscopy gave support to this finding. The maximum in electrical resistivity and minimum in dielectric losses arising from heat-treatment at
20.
UP
o 10 "5
.'r
t
600
t
[
7OO 8O0 9O0 Heat treatment temperature (°C.)
4
8 12 16 "lime(hours)
20
24
Fig. 3. The effects of heat-treatment temperature on the power factor of a glass-ceramic; measurements made at 300°C and a frequency of 100 kHz. Fig. 4. Heat evolution during isothermal heat-treatment at 730°C for MgO-AI2Oa-B203 glassceramic. A, nucleated at 6 7 0 ° C / 4 h; B, nucleated at 670°C/12 h.
P.W. McMillan / The crystallisation of glasses
75
750°C result from the incorporation of lithium ions into crystalline lithium disilicate in which their mobility is greatly reduced as compared with that in the glass phase. These results and those from similar investigations underline the value of electrical measurements as a sensitive tool for the study of microstructural changes brought about by glass crystallisation. Another valuable class of investigative techniques is represented by thermal analytical methods. These include differential thermal analysis (DTA) and differential scanning calorimetry (DSC). The use of these techniques can be illustrated by the results of Lau and McMillan [19] who investigated the crystallisation of a glass of the MgO-A1203-B203 type. Preliminary trials indicated that conversion of this material to a fine-grained glass-ceramic could be achieved by a nucleation treatment at 670°C and a crystallisation treatment at 730°C. Using a differential scanning calorimeter, the rate of heat evolution at 730°C was determined for specimens that had received various nucleation treatments. Typical results are shown in fig. 4. A striking feature of the results is the change in slope of the curves after heat-treatment at 730°C for about 4 h. It was also noted that the glass-ceramics produced were transparent for heat-treatments up to 4 h but became opaque thereafter. The change in slope of the DSC curves was taken to indicate that two crystalline phases involving
2"
i/
..2"
~0 i5 i-2-
-4" A, / S 600
700
800
Temperature ('C)
900
9 10 Int.(sec.)
1
Fig. 5. DTA curves for MgO-AI203-B203 glass ceramic. A, nucleated at 670°C/12 h; B, nucleated at 670°C/4 h; C, no nucleation treatment. Fig. 6. DSC data for MgO-AI203-B203 glass-ceramic. A, nucleated at 670°C/4 h; B, nucleated at 670°C/12 h.
P. W. McMillan / The crystallisation of glasses
76
different kinetics were being formed. DTA studies, illustrated in fig. 5, confirmed that for specimens that had received a nucleation treatment, two crystalline phases were being formed as shown by the occurrence of two exothermic peaks. It was deduced that the lower temperature peak related to the crystal phase responsible for the initial portion of the DSC curve. Further DTA studies showed that the higher temperature peak was shifted to lower temperatures by the use of finer glass powder whereas the lower peak was not affected. This result indicated that the higher temperature phase was surfacenucleated while the other phase was bulk-nucleated. Further analysis of the DSC data confirmed this finding. From plots of the type illustrated in fig. 6 it is possible to derive the Avrami exponent, n, from the slopes of the two straight line regions. It was found that for the lower temperature phase the value of n was 3 + 0.2. This implies a constant growth rate and a nucleation rate of zero at the treatment temperature of 730°C. For the higher temperature phase, the value of n was found to lie between 0.6 and 1, which could imply either the thickening of plates when complete edge impingement has occurred or the growth of needle-shaped crystals. Either possibility is consistent with surface-nucleated crystallisation.
References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13] [14] [15] [16] [17] [18] [19]
W.H. Zachariasen, J. Am. Ceram. Soc. 54 (1932) 3841. H. Moore and P.W. McMillan, J. Soc. Glass Technol. 40 (1956) 66. H.L. Trap and J.M. Stevels, Glastech. Ber. 32K (1959) 51. D. Turnbull and M.H. Cohen, J. Chem. Phys. 29 (1958) 1049. R.D. Maurer, J. Appl. Phys. 33 (1962) 2132. P.F. James and P.W. McMillan, Phil. Mag. 18 (1968). M. Tomozawa, in: Advances in Nucleation and Crystallization in Glasses (American Ceramic Society, 1971) p. 41. K. Matusita, S. Sakka, T. Muki and M. Tashiro, J. Mat. Sci. 10 (1975) 94. P.W. McMillan, Proc. Xth Int. Cong. on Glass 14 (1974) 1. P.W. McMillan and S.V. Phillips, in: Glass-Ceramics (Academic Press, London, 1979) p. 53. S.V. Phillips and P.W. McMillan, Glass Technol. 7 (1965) 121. B. Hibberd and P.W. McMillan, in: Glass-Ceramics (Academic Press, London, 1979) p. 54. A. Chlebik and P.W. McMillan, J. Non-Crystalline Solids 38 (1980) 509. J.S. Olcott and S.D. Stookey, in: Advances in Glass Technology (Plenum, New York, 1962) p. 400. G. Partridge and P.W. McMillan, Glass Technol. 15 (1974) 127. R. Adams and P.W. McMillan, J. Mat. Sci. 17 (1982) 2727. P.W. McMillan, Glass-ceramics (Academic Press, London, 1979). D.I.H. Atkinson and P.W. McMillan, J. Mat. Sci. 11 (1976) 994. J. Lau and P.W. McMillan, to be published.