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Effect of fining with sodium chloride on the phase separation of a soda lime silica glass
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Journal of Non-Crystalline Solids 208 (1996) 294-302
Effect of fining with sodium chloride on the phase separation of a soda lime silica glass A. Hoell a,*, R. Kranold a, U. Lembke a, j. Aures b,1 a Universitk~t Rostock, Fachbereich Physik, Universiti~tsplatz 3, D-18051 Rostock, Germany b Jenaer Glaswerk GmbH, Otto-Schott-Strafle 13, D-07745 Jena, Germany Received 15 November 1995; revised 1 April 1996
Abstract
The effect of the technologically important procedure of fining with NaC1 on the liquid-liquid phase separation of the 13Na20-11CaO-76SiO 2 (mol%) glass in the metastable region of the miscibility gap has been studied by measuring the progressive changes of the small angle X-ray scattering (SAXS) intensities that characterize the phase separation process at a constant temperature of 600°C. It has proved that a small quantity of chloride remains in the glass after fining. This small content of chloride causes a dramatic change of kinetics and equilibrium conditions for the phase separation process in the glass in comparison with a pure soda lime silica glass: the phase separation is accelerated considerably and the maximum volume fraction of the silica droplet phase is increased. The effect of chloride is interpreted as a shift of the miscibility gap position by 45 K up to higher temperatures. This explanation is confirmed by measurements of the binodal temperatures of the pure glass and the glass contaminated with chloride. Viscosity measurements show that chloride does not significantly alter the conditions for diffusion in the glass. Thus, the effect of chloride on the phase separation is mainly attributed to a rise of the miscibility gap.
I. Introduction
Liquid-liquid phase separation is a widespread phenomenon in silicate glasses [1-3]. Since phase separation has an effect on the physical properties and the crystallization behaviour of glasses, a detailed understanding of this process can provide a basis for the controlled manufacture of numerous glassy materials, as, for example, opalescent glasses
* Corresponding author. E-mail: [email protected]. 1 Present address: NetConsult GmbH EbertstraSe 16, D-07743 Jena, Germany.
and glass ceramics. On the other hand, phase separation must be avoided for most optical glasses. The thermodynamic conditions for phase separation in glasses are mainly determined by the batch composition and the course of the heat treatment. A multitude of influences on the tendency for phase separation has to be considered during the technological process of glass preparation: loss in concentration of certain components during melting, melting atmosphere, fining agents and fining duration [4], thermal treatment to avoid mechanical stresses, and rates of cooling down to room temperature. The aim of the present paper is to investigate one particular topic: the effect of small amounts of chloride ions on the phase separation in a soda lime silica glass.
0022-3093/96/$15.00 Copyright © 1996 Elsevier Science B.V. All rights reserved. PII S 0 0 2 2 - 3 0 9 3 ( 9 6 ) 0 0 5 1 5 - 7
A. Hoell et al. / Journal of Non-Crystalline Solids 208 (1996) 294-302
These CI- impurities are shown to remain in the glass after fining with sodium chloride. Up to now, C1- impurities have been assumed to have a negligible influence on phase separation. In previous papers, a remarkable influence on the phase separation in sodium silicate glasses has been attributed to traces of O H - and F - impurities mainly [5-13]. Weyl and Marboe [12] pointed out that water is perhaps the most powerful nucleation catalyst for silicates. It acts in a manner similar to that of fluorine. In silicate glasses, O H - ions [8] and F ions [1] lead to a rupture of the strong S i - O - S i bond. This rupture is the reason for a considerable decrease of the viscosity of the glass. Increased coarsening rates were found with increasing O H content in sodium silicate glasses [6,7]. Furthermore F - causes an increase of the miscibility gap [5]. Contrary to the detailed results which have been reported on the influence of O H - and F - ions on the immiscibility of silicate glasses, only a few qualitative results concerning the role of chloride ions were published [9-11]. Essentially, the effect of C1- in the glass is discussed by analogy with the action of fluorine ions [10], without investigating it in detail. This situation of a vague knowledge about the effect of chloride arises mainly due to the small solubility of chloride in silicate glasses and, consequently, the limited use of chloride as a glass component [10]. However, fining with sodium chloride is a method frequently used to homogenize silicate melts. A glass of the composition 13Na20-11CaO76SIO 2 (mol%) was chosen because it is thoroughly investigated [14-19]. This composition is near the boundary of the well known metastable immiscibility gap. According to a pseudobinary tie line the glass separates into two phases, silica enriched droplets and matrix. The phase separation proceeds via a binodal mechanism. That means that the process has three stages which are nucleation, growth, and ripening [11]. Small angle X-ray scattering (SAXS) [20,21] has proved to be useful for the characterization of submicroscopic heterogeneities in glasses [3,22]. The method provides a variety of structural parameters which represent average values obtained from a set of billions of particles. By evaluating the evolution of these structural parameters the kinetics of the phase separation process can be deduced [22].
295
2. Experimental 2.1. Preparation and characterization of glass samples Two glasses, denoted A and D, of composition 13Na20-11CaO-76SiO 2 (mol%), were prepared from c.p. grade SiO 2 powder (fine grain size distribution; content of ferric oxide < 14 ppm), Na2CO 3, NaNO 3, and CaCO 3. To the batch, prepared to produce 1700 g of glass A, 17 g of the fining agent NaC1 were added. The two batches were treated identically during melting. They were melted at 1530°C in air for 13.5 h. A closed 100% platinum crucible and an electric furnace with Superkanthal heating elements were used. After 4 h the liquids were stirred with a platinum stirrer with 10 to 40 rpm for 9 h to ensure maximum homogeneity. The liquids were poured into steel moulds coated with A1203 powder, which were preheated at 580°C. Starting at this temperature, which is just above the glass transition temperature, Tg, the glasses were cooled to room temperature with cooling rates decreasing from 4 K min-1 to 1 K min-1 in order to avoid mechanical stress. The glass blocks produced in this manner were free of bubbles and flaws and did not show any sedimentation effect. By means of wet chemical analysis, only small differences between the two glass compositions were found. A Na20 deficit of about 0.8 mol% with regard to the batch composition has been established for the final glass compositions (Table 1). The CaO content in glass A fits the batch composition, whereas the CaO content in glass D is 0.5 mol% greater. The main difference is a chloride content of 0.47 mol% in glass A due to the fining agent that partially remains in the glass melt. The glasses have been analyzed thoroughly in order to detect other impurities. The results of atomic absorption spectroscopy proved that the contents of different kinds of impurities amount to a few ppm and do not differ significantly in the glasses investigated. Using X-ray absorption spectroscopy the same result was obtained for the contamination of the glass melt with platinum of the melting crucibles. The water contents are very small in the glasses A and D, as revealed by infrared absorption spectroscopy. These measurements were performed with polished glass platelets of 1.5 mm
A. Hoell et aI. / Journal of Non-Crystalline Solids 208 (1996) 294-302
296
thickness. The detection limit for OH- was 10 ppm. Since no OH- signal was registered for either sampies A and D, their hydroxyl content should be lower than a few ppm. The glass transition temperatures, Tg, the densities, d, and the viscosities, r/, of the as-prepared glasses are listed in Table 1. The Tg'S were determined by means of differential thermal analysis with an uncertainty of ___5 K. The measurements of the sample densities were carried out in toluol at 20°C by means of the Archimedes principle. The viscosities were measured at (600 + 1)°C with the fiber elongation method. The measurements were performed with rods of dimension 1 × 1 × 51 mm 3. The binodal temperature, which characterizes the boundary between the metastable and the homogeneous region in the quasi binary phase diagram for a given glass composition, was determined with the opalescence method [15]. Glass strips, about 10 cm long, were heat treated beneath the binodal temperature for a time long enough to ensure the development of an appreciable phase separation throughout the samples. Afterwards the samples were placed in a furnace with a defined, nearly linear thermal gradient. After air quenching to room temperature the glass bars appeared to be homogeneous at the high temperature end and remained phase separated in the temperature region of the miscibility gap below the binodal temperature. The latter is indicated by a strong light scattering of the low temperature side of the glass bar. The binodal temperature corresponds to the position on the bar where the region of maximum opalescence is demarcated from the clear homogeneous glass. Calibrating the temperature field within the furnace this position can be assigned to the corresponding critical temperature with an uncertainty of + 10 K. In order to produce specimens for the SAXS
investigations, the blocks of the glasses A and D were cut into samples of about 25 mm × 10 mm × 3 mm. The as-prepared samples were heat treated isothermally at 600°C for different times. The heat treated samples did not show any crystallization effect. The samples were ground to a final thickness of 0.1 mm and polished using cerium oxide. The surface roughness obtained corresponds to optical quality in order to avoid surface scattering of X-rays. 2.2. SAXS measurements
The X-ray source was a sealed tube with copper anode operating at power settings of 34 kV and 40 mA. Monochromatic Cu K~ radiation was obtained using a Ni filter in combination with a proportional counter with pulse height analyzer. The scattering intensities, l(s), were measured between Smin 0.09 nm -1 and Smax = 3 . 0 nm -1 with a commercial Kratky camera using infinitely long slit collimation conditions [21]. The quantity s denotes the modulus of the scattering vector, s = 41r s i n ( 0 / 2 ) / h , where h is the X-ray wavelength and 0 is the scattering angle. The scattering curve recorded under these conditions contains the complete information about all scattering particles in the size range between 1 nm and 40 nm. All the scattering curves were measured in several runs. The statistical error was less than 3% for any measuring point. After correcting for background scattering the recorded scattering intensities were converted into electron units by means of a calibrated polyethylene (Lupolen) standard sample [21]. From the scattering curves, I(s), the following structural parameters are calculated: the average particle radius, R~, the particle number density, Nv, the internal surface per unit volume, Sv, and the mean square of electron density fluctuations, ((A p)2 ). "------
Table 1 Chemical and physical data of the glasses A and D Glass
A D Batch
Composition (mol%) Na20
CaO
SiO 2
CI
12.13 12.20 13.00
11.02 11.48 11.00
76.38 76.32 76.00
0.47 -
Tg (°C)
Density, d ( g / c m 3)
Viscosity at 600°C, r/(Pa s)
569 570
2.458 2.472
4.0 X l0 II 4.0 × 1011
A. Hoell et al. / Journal of Non-CrystallineSolids 208 (1996)294-302 The mean square of electron density fluctuations, ((A p)2 ), is determined from the first moment of the scattering curve, I(s), by the expression [20] 1
( ( A p ) 2 ) - 2~Aa fo s l ( s ) ds,
(1)
where the quantity a is the distance between the specimen and the detector slit. The scattering intensities recorded between Smin and Smax are extrapolated to zero angle using the Guinier law [20] and to infinity using the Porod law [20]. The Porod law for infinitely long slit collimation is I(s) = k/s 3
(2)
for large values of s. The constant k is connected with the internal surface per unit volume, S,., via
~: = ( ~ / 2 ) A a ( Ap)2S,,.
(3)
For a two-phase system the mean square of electron density fluctuations can be calculated as [20] ( ( A p ) 2 > = w(1 - w ) ( A p ) 2,
(4)
where w and (1 - w) denote the volume fractions of the droplet phase and the matrix phase, respectively, and
A p = UA[(dz/M)droptet- (dz/M)matrix]
R,
foRaN(R)dR
3 foS[(s) ds - f o R 3 N ( R ) d g "
size distribution. It has been demonstrated elsewhere [24] that R t is suitable to test certain power laws predicted by different theories for particle growth during a phase separation process. The particle number density, N,., has been calculated using Eqs. (3) and (6):
St N,
47rR~ "
(7)
3. R e s u l t s
The results of the SAXS investigations indicate considerable differences in the course of the phase separation process in the glasses A and D. This feature is demonstrated in Fig. 1 for the scattering curves of the specimens as-prepared and heat treated at 600°C for 5.5 h. While the untreated sample of glass A already shows an advanced stage of phase separation indicated by an intense small angle scattering effect, no effect due to phase separation is observed in the flat SAXS curve of the untreated sample of glass D. With progressing heat treatment the SAXS effect develops and intensifies. In Fig. 2, the evolution of the structural parame-
(5)
is the difference between the electron densities of the two phases. NA is Avogadro's number, d is the density, z and M are the number of electrons and the molar weight per unit of composition of the two phases. In the present case of the phase separated glass the droplets are enriched with silica at the expense of the silica concentration in the matrix glass. Because of the nearly spherical shape of the droplets the average radius, R 1, which can be extracted from the SAXS data with high accuracy [23], is an appropriate structural parameter for characterizing the particle size. RI can be calculated from the intensities, I(s), according to [20,24]
4 f o I ( s ) ds
297
(6)
As can be seen from Eq. (6), R 1 represents the mean radius of the particle-volume distribution of the droplet phase [23]. N(R) is the corresponding
1027
A (5.5 h) 1026
D (5.5 h) A(0h)
1025 D (0 h)
1024 10-2
~
i
IiA,,,I
. . . .
,,,,I
10o s=4~sin(0/2)/?~ 10-1
,
,
,,,,,,
10 ] Into q]
Fig. 1. SAXS curves of the specimensof glass A (with Cl- ) and glass D (without C1- ) as-prepared and heat treated at 600°C for 5.5 h.
298
A. Hoell et al. / Journal of Non-Crystalline Solids 208 (1996) 294-302
lO
20
101 W"
10 0 5.10 5
0
r-------m
~ g
10 5
i
r
i
ilrlll
,
. . . . . . .
till
r
i
1017
T
I
i
i
,~tl
i
,
,
,,,i,I
10
1
D
I
0
100 t [h]
Fig. 3. Mean square of electron density fluctuations, ((Ap)2), and corresponding volume fraction, w, of the nucleus phase in dependence on the time, t, of the isothermal heat treatment. The relative precisions were estimated to be 3% in ((A p)2 ), which do not exceed the extension of the points. The solid lines are drawn as guides for the eye. The value of win,x (dashed line) is calculated from Eq. (8).
.o.
Z
1016
1015
.................. 10
1
100 t [h]
Fig. 2. Average particle radius, R 1, internal surface per unit volume, Sv, and particle number density, N~., of glass A (with C1- ) and glass D (without CI- ) in dependence on the time, t, of the isothermal heat treatment at 600°C. According to [23] the relative precisions were estimated to be 3% in R 1, 8% in S,,, and 14% in N~,. In the case of R], the relative precisions do not exceed the extension of the points.
dashed line in Fig. 3 indicates the theoretical maxim u m value of ( ( A p ) 2) for the case o f complete phase separation, which was calculated using Eqs. (4) and (5) taking into account the metastable i m m i s cibility boundaries experimentally determined b y H a m m e l [15] for the quasi b i n a r y system SiO 2 [ ( 1 3 / 2 4 ) N a 2 0 - ( 1 1 / 2 4 ) C A O ] ( m o l % ) shown in Fig. 4. Here, the m a x i m u m v o l u m e fraction, Wmax , obtain-
1200 o~
ters R 1, S v, and Nv during phase separation at 600°C is shown for glass A containing chloride and glass D without any chloride. O w i n g to fining with 1 wt% NaC1 an appreciable effect on the phase separation behaviour of the glass investigated has been observed. As demonstrated in Fig. 2 a considerable increase of both the particle size and the growth velocity is obtained due to the chloride addition into the melt of glass A. Furthermore, the interface area, S~., and the particle n u m b e r density, N,,, in glass A are larger and they decrease more rapidly than in glass D. Fig. 3 shows the evolution of the m e a n square of electron density fluctuations, ( ( A p ) 2 ) , determined from the SAXS data according to Eq. (1). The
1000 800
o
600 t 400 I I
60
70
f
80
90
100
Xsio= [tool%] Fig. 4. Miscibility gap in the soda lime silica glass according to the quasibinary tie line (data given by Hammel [15]), SiO2[(13/24)NazO-(ll/24)CaO] (mol%). For T=600°C the silica concentrations of the homogeneous glass, Xg = 0.76, of the nuclei, x, = 0.99, and of the matrix, x m = 0.73, are marked. The dashed line indicates the shift of x m and x, caused by a rise of the miscibility gap by 45 K up to higher temperatures.
A. Hoell et al. / Journal of Non-Crystalline Solids 208 (1996) 294-302
able for the glass without any impurity can be calculated for a given treatment temperature from the equilibrium mole fractions of the silica enriched nuclei, x,, and of the matrix, x m, according to the known miscibility gap and the composition of the initial homogeneous glass, Xg, (Xg - x m ) d g M n
(8)
299
proceeds with ripening. Investigating the time dependence of the structural parameters during heat treatment, SAXS offers the opportunity to distinguish between different stages of the phase separation process [3,22,26]. For the chloride containing glass A the following power laws are deduced from Fig. 2:
N, OCt - l ' 1 8 ;
g 1 ~ t0"37;
M and d are the molar weights and the mass densities [25] of the phases denoted by the corresponding indices. Using the experimental data for the miscibility gap given by Hammel [15] the following values are obtained for T = 600°C: Xg = 0.76, x m = 0.73, and x n = 0.99 which are marked in Fig. 4. The density of the initial homogeneous glass was measured to be dg = 2.48 g cm -3. For the purpose of calculating an upper limit for Wmax and for the mean square of electron density fluctuations, ((A p)2), it was assumed that the nuclei consist of pure silica. The corresponding mass density is d n = 2.2 g cm -3. With these data and the Eqs. (4), (5), and (8) the theoretical values of Wmax = 0.13 and ((A p)2 )calc = 7.2 × 1044 cm -6 are obtained and marked by the dashed line in Fig. 3. This figure demonstrates that for long treatment times the experimental ((A p)2) values of the pure glass D should fit asymptotically the calculated value of 7.2 × 1044 cm -6 whereas for glass A, that contains chloride, the equilibrium value of ( ( A p ) 2) is considerably larger and can be estimated to amount to at least 8.9 × 1044 cm -6. Different binodal temperatures were determined for the glasses A and D. The binodal temperature of glass A, Thin,A, amounts to (770 + 10)°C whereas the binodal temperature of glass D, Thin,D, was obtained to be (705 ___10)°C.
If these results are compared with the theoretical predictions of the classical theory of Lifshitz and Slyozov [27] and Wagner [28] (LSW theory) for the asymptotic stage of diffusion-limited ripening,
4. Discussion
4.1. Kinetics The compositions of the glasses A and D lie within the metastable region of the phase diagram. That means that phase separation is going on according to the binodal mechanism, that is characterized by nucleation and growth, which in the late stage
(R)~tl/3;
N,.~t-1;
St, ~ t -0"37.
(9)
Wmax= ( X - Xm)d, Mg
S~,at -1/3,
(10)
it can be concluded that glass A has already reached the asymptotic stage of ripening. The appreciable volume fraction of the nucleus phase, w, which is increasing up to 16.7% in glass A, has no effect on the time exponents in the ripening equations, but has only an influence on the prefactors in the predicted power laws and on the form of the size distribution function [29]. This fact is important because the LSW theory is valid for an infinitely small volume fraction of the droplet phase only. The phase separation kinetics of the pure glass D, in relation to glass A, are considerably slower (Fig. 2), which results in R 1 (xtl/4;
g ~ . ~ t -0"48
S ~ ( x t -0"02 .
(11)
The particle number density, N~,, decreases with progressing treatment time, t, whereas the internal surface, S,., of the scattering system decreases only slightly. It is worth mentioning that in the examined time interval of heat treatment the volume fraction of the precipitating nuclei, w, is less than the saturation value, Wmax (Fig. 3). The power law exponents in (11) for the investigated structural parameters of glass D (Fig. 2) are not typical, neither for the asymptotic stage of diffusion-limited ripening nor for the stage of independent growth. A similar behaviour has already been observed experimentally for a soda lime silica glass investigated by Burnett and Douglas [ 18]. They found a so-called intermediate stage between diffusionlimited growth and the asymptotic stage of
300
A. Hoell et al. / Journal of Non-Crystalline Solids 208 (1996) 294-302
diffusion-limited ripening. This intermediate stage has a retardation of the increase of the mean particle radius ( R ) during thermal treatment leading to a growth velocity which is temporarily even smaller than in the following stage of asymptotic ripening. Theoretical computations that model the evolution of the nucleus size distribution and the mean particle radius during nucleation, growth, and ripening in viscoelastic media [30] suggest that this phenomenon is typical for phase separating glasses where the relative supersaturation of the initial homogeneous glass in the metastable region of the miscibility gap is relatively large. Apparently the chloride content in glass A acts in such a manner that the phase separation process is accelerated in comparison to glass D which has the same composition ratio concerning the oxide components but does not contain chloride. Therefore, the main result of the kinetic studies is the fact that the chloride containing glass A has already reached the asymptotic stage of ripening whereas for the same treatment conditions glass D is in the intermediate growth stage. Consequently, in case of glass A, the volume fraction and the magnitude of the mean square of electron density fluctuations after the 100 h heat treatment should match with the corresponding saturation values which are valid for the equilibrium case. However, in glass D (without chloride) the maximum phase volume, Win,x (dashed line in Fig. 3), does not precipitate even during a 100 h heat treatment.
4.2. Miscibility gap As demonstrated in Fig. 3, the experimentally obtained magnitude of ( ( A R ) 2) = 8.9 X 1044 c m - 6 for glass A exceeds the theoretical value of 7.2 x 1044 cm -6 calculated from the miscibility gap data of the pure soda lime silica glass by 23%. The influence of the small differences between the compositions of the glasses A and D and, therefore, slight changes of the initial composition in the miscibility gap turns out to correspond to a change o f ( ( A p ) 2) not exceeding +5%. This conclusion was confirmed in Ref. [26]. Obviously, the chloride content in glass A is responsible for the deviation of the final state of phase separation from that expected for the pure
glass. According to Eq. (4) two possible reasons for this effect must be discussed. On the one hand an increased contrast, A p, between the electron densities of the two phases could be caused by the chloride content and on the other hand a higher maximum volume fraction could be developed. Analyzing first a possible change of A/9 one has to consider that the electron density of chloride is very similar to that of the other glass components. The quantitative changes of A p which are obtainable by considering the effect of the chloride content in glass A were estimated to be at least one order of magnitude smaller than necessary in order to explain the increase of ((Ap) 2) observed experimentally. Consequently, the experimental result has to be discussed in terms of an increasing precipitated phase volume, that is larger than w calculated by Eq. (8), owing to the chloride impurity of the glass composition and resulting in a parallel shift of the immiscibility boundary to higher temperatures as illustrated by the dashed binodal curve in Fig. 4. The parallel shift is justified by the results of Kawamoto and Tomozawa [31] who found that the binodal curves of the alkali and earth alkaline silicate glasses can be reduced to one normalized binodal master curve. Furthermore, Kawamoto et al. showed in [32] that the binodal curves of the NazO-B203-SiO 2 system with small amounts of MoO 3 can be represented by curves parallel to the binodal curve of the pure glass. According to Eqs. (4) and (8) the experimentally observed increase of ((Ap) z) by 23% caused by the chloride addition in glass A corresponds to a gap shift of 45 K up to higher temperatures. This result is qualitatively confirmed by the direct measurements of the binodal temperatures of glass A and glass D using a gradient furnace. The thermodynamic driving force, A/z, of a phase separation process depends on the differences in the position of the miscibility gap and the silica concentration of the initial homogeneous glass, Xg, at the temperature of heat treatment, T. Consequently, the increase of the binodal temperature, Tbin, in glass A leads to an increase of A/x which, in a first approximation, is proportional to the undercooling, Tbin -- T [15,26]. Accordingly, the parallel shift of the immiscibility boundary in the quasi binary phase diagram should result in modified kinetic parameters for the whole phase separation process.
A. Hoell et al. / Journal of Non-Crystalline Solids 208 (1996) 294-302
The rate of homogeneous nucleation, given by [33] Inud=loexp(-[AG*
+ AGo]/kT )
/nucl'
is
(12)
with the activation energy for the formation of a critical spherical nucleus 16 cr 3 AG* = - - ~ r 3 A~ 2 "
(13)
301
nucleation and growth in glass A is stimulated by the increased A/x according to Eqs. (13) and (14), whereas an influence of the remaining chloride on the diffusion coefficient, D*, can be to a great extent excluded. Additionally, the higher thermodynamic driving force in glass A results in higher nucleation rates and growth rates (Eqs. (12)-(14)). Consequently, at the end of the stage of independent growth glass A contains more and larger nuclei than glass D (Fig. 2).
The power law for diffusion-limited growth reads [33] (R) 2 =
2D*A/zc~
5. Conclusions t,
(14)
ckT
where k is the Boltzmann constant, T the absolute temperature, AG o the activation energy of diffusion, or the interface tension, D* the effective diffusion coefficient, c= the concentration of the diffusing species in the undisturbed matrix at infinite distance from the separated particle, and c the concentration of the diffusing species in the nucleus. From Eqs. (12) and (13) it can be derived that a rising thermodynamic driving force increases the nucleation rate provided that the activation energy of diffusion, A G o , remains constant. Taking into account the proportionality between the diffusion coefficient, D*, and the reciprocal viscosity of the glass [1], the invariance of A G o can be tested by means of viscosity measurements. For the glasses A and D the same viscosity, 7/= 4.0 × 1011 Pa s, was obtained at 600°C. That means that D* and A G o should have approximately the same values in both glasses. The measured value for matches that determined by Hammel et al. [16] for a 13Na20-11CaO-76SiO 2 glass. These findings correspond to the experiences made by different authors, to which reference is made in Ref. [34], that for soda lime silica glasses in the composition range investigated the viscosity depends only slightly on composition. The identical values of the glass viscosities evidently prove that the differences observed for the phase separation kinetics of the glasses A and D are mainly due to a changed thermodynamic driving force, A/x. The chloride that remained in glass A after the fining is the reason for an increased A/x in comparison to glass D. The observed acceleration of
Small quantities of chloride were proved to remain in the glass melt after fining of soda lime silica glass with sodium chloride. The fined glass A investigated contains 0.47 mol% chloride. This small content of chloride causes changes of kinetics and equilibrium conditions for the phase separation process in the glass in comparison with a pure soda lime silica glass: the phase separation is accelerated and the maximum volume fraction of the droplet silica phase is increased. The effect of chloride is interpreted by a shift of the miscibility gap position to higher temperatures of about 45 K. Viscosity measurements have shown that chloride does not significantly alter the conditions for diffusion in the glass. Thus, the effect of chloride on phase separation is a promotion of the phase separation process caused by a rise of the miscibility gap, which increases the thermodynamic driving force.
Acknowledgements Financial support of the Deutsche Forschungsgemeinschaft (grant Kr 137212-1) is gratefully acknowledged.
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[19] K. Huntebrinker, U. Wtirz and G.H. Frischat, J. Non-Cryst. Solids 11 (1989) 189. [20] A. Guinier and G. Fournet, Small Angle Scattering of X-rays (Wiley, New York, 1955). [21] O. Glatter and O. Kratky, Small Angle X-ray Scattering (Academic Press, London, 1982). [22] R. Kranold and G. Walter, in: Physical Research 12: Amorphous Structures - - Methods and Results, ed. D. Schulze (Akademie Verlag, Berlin, 1990) p. 267. [23] G. Walter, R. Kranold, T. Gerber, J. Baldrian and M. Steinhart, J. Appl. Cryst. 18 (1985) 205. [24] J. MiSller, R. Kranold, J. Schmelzer and U. Lembke, J. Appl. Crystallogr. 28 (1995) 553. [25] R.R. Shaw and D.R. Uhlmann, J. Non-Cryst. Solids 1 (1969) 474. [26] A. Hoell, PhD thesis, Universit~it Rostock (1991). [27] J.M. Lifshitz and V.V. Slyozov, J. Phys. Chem. Solids 19 (1961) 35. [28] C. Wagner, Z. Elektrochem. 65 (1961) 581. [29] J.H. Yao, K.R. Elder, H. Guo and M. Grant, Phys. Rev. B47 (1993) 14110. [30] J. Bartels, U. Lembke, R. Paskova, J. Schmelzer and I. Gutzow, J. Non-Cryst. Solids 136 (1991) 181. [31] Y. Kawamoto and M. Tomozawa, Phys. Chem. Glasses 22 (1981) 11. [32] Y. Kawamoto, K. Clemens and M. Tomozawa, J. Am. Ceram. Soc. 64 (1981) 292. [33] I. Gutzow and J. Schmelzer, The Vitreous State (Springer, Berlin, 1995). [34] B. Kumar and G.E. Rindone, Phys. Chem. Glasses 20 (1979) 119.
ELSEVIER
Journal of Non-Crystalline Solids 208 (1996) 294-302
Effect of fining with sodium chloride on the phase separation of a soda lime silica glass A. Hoell a,*, R. Kranold a, U. Lembke a, j. Aures b,1 a Universitk~t Rostock, Fachbereich Physik, Universiti~tsplatz 3, D-18051 Rostock, Germany b Jenaer Glaswerk GmbH, Otto-Schott-Strafle 13, D-07745 Jena, Germany Received 15 November 1995; revised 1 April 1996
Abstract
The effect of the technologically important procedure of fining with NaC1 on the liquid-liquid phase separation of the 13Na20-11CaO-76SiO 2 (mol%) glass in the metastable region of the miscibility gap has been studied by measuring the progressive changes of the small angle X-ray scattering (SAXS) intensities that characterize the phase separation process at a constant temperature of 600°C. It has proved that a small quantity of chloride remains in the glass after fining. This small content of chloride causes a dramatic change of kinetics and equilibrium conditions for the phase separation process in the glass in comparison with a pure soda lime silica glass: the phase separation is accelerated considerably and the maximum volume fraction of the silica droplet phase is increased. The effect of chloride is interpreted as a shift of the miscibility gap position by 45 K up to higher temperatures. This explanation is confirmed by measurements of the binodal temperatures of the pure glass and the glass contaminated with chloride. Viscosity measurements show that chloride does not significantly alter the conditions for diffusion in the glass. Thus, the effect of chloride on the phase separation is mainly attributed to a rise of the miscibility gap.
I. Introduction
Liquid-liquid phase separation is a widespread phenomenon in silicate glasses [1-3]. Since phase separation has an effect on the physical properties and the crystallization behaviour of glasses, a detailed understanding of this process can provide a basis for the controlled manufacture of numerous glassy materials, as, for example, opalescent glasses
* Corresponding author. E-mail: [email protected]. 1 Present address: NetConsult GmbH EbertstraSe 16, D-07743 Jena, Germany.
and glass ceramics. On the other hand, phase separation must be avoided for most optical glasses. The thermodynamic conditions for phase separation in glasses are mainly determined by the batch composition and the course of the heat treatment. A multitude of influences on the tendency for phase separation has to be considered during the technological process of glass preparation: loss in concentration of certain components during melting, melting atmosphere, fining agents and fining duration [4], thermal treatment to avoid mechanical stresses, and rates of cooling down to room temperature. The aim of the present paper is to investigate one particular topic: the effect of small amounts of chloride ions on the phase separation in a soda lime silica glass.
0022-3093/96/$15.00 Copyright © 1996 Elsevier Science B.V. All rights reserved. PII S 0 0 2 2 - 3 0 9 3 ( 9 6 ) 0 0 5 1 5 - 7
A. Hoell et al. / Journal of Non-Crystalline Solids 208 (1996) 294-302
These CI- impurities are shown to remain in the glass after fining with sodium chloride. Up to now, C1- impurities have been assumed to have a negligible influence on phase separation. In previous papers, a remarkable influence on the phase separation in sodium silicate glasses has been attributed to traces of O H - and F - impurities mainly [5-13]. Weyl and Marboe [12] pointed out that water is perhaps the most powerful nucleation catalyst for silicates. It acts in a manner similar to that of fluorine. In silicate glasses, O H - ions [8] and F ions [1] lead to a rupture of the strong S i - O - S i bond. This rupture is the reason for a considerable decrease of the viscosity of the glass. Increased coarsening rates were found with increasing O H content in sodium silicate glasses [6,7]. Furthermore F - causes an increase of the miscibility gap [5]. Contrary to the detailed results which have been reported on the influence of O H - and F - ions on the immiscibility of silicate glasses, only a few qualitative results concerning the role of chloride ions were published [9-11]. Essentially, the effect of C1- in the glass is discussed by analogy with the action of fluorine ions [10], without investigating it in detail. This situation of a vague knowledge about the effect of chloride arises mainly due to the small solubility of chloride in silicate glasses and, consequently, the limited use of chloride as a glass component [10]. However, fining with sodium chloride is a method frequently used to homogenize silicate melts. A glass of the composition 13Na20-11CaO76SIO 2 (mol%) was chosen because it is thoroughly investigated [14-19]. This composition is near the boundary of the well known metastable immiscibility gap. According to a pseudobinary tie line the glass separates into two phases, silica enriched droplets and matrix. The phase separation proceeds via a binodal mechanism. That means that the process has three stages which are nucleation, growth, and ripening [11]. Small angle X-ray scattering (SAXS) [20,21] has proved to be useful for the characterization of submicroscopic heterogeneities in glasses [3,22]. The method provides a variety of structural parameters which represent average values obtained from a set of billions of particles. By evaluating the evolution of these structural parameters the kinetics of the phase separation process can be deduced [22].
295
2. Experimental 2.1. Preparation and characterization of glass samples Two glasses, denoted A and D, of composition 13Na20-11CaO-76SiO 2 (mol%), were prepared from c.p. grade SiO 2 powder (fine grain size distribution; content of ferric oxide < 14 ppm), Na2CO 3, NaNO 3, and CaCO 3. To the batch, prepared to produce 1700 g of glass A, 17 g of the fining agent NaC1 were added. The two batches were treated identically during melting. They were melted at 1530°C in air for 13.5 h. A closed 100% platinum crucible and an electric furnace with Superkanthal heating elements were used. After 4 h the liquids were stirred with a platinum stirrer with 10 to 40 rpm for 9 h to ensure maximum homogeneity. The liquids were poured into steel moulds coated with A1203 powder, which were preheated at 580°C. Starting at this temperature, which is just above the glass transition temperature, Tg, the glasses were cooled to room temperature with cooling rates decreasing from 4 K min-1 to 1 K min-1 in order to avoid mechanical stress. The glass blocks produced in this manner were free of bubbles and flaws and did not show any sedimentation effect. By means of wet chemical analysis, only small differences between the two glass compositions were found. A Na20 deficit of about 0.8 mol% with regard to the batch composition has been established for the final glass compositions (Table 1). The CaO content in glass A fits the batch composition, whereas the CaO content in glass D is 0.5 mol% greater. The main difference is a chloride content of 0.47 mol% in glass A due to the fining agent that partially remains in the glass melt. The glasses have been analyzed thoroughly in order to detect other impurities. The results of atomic absorption spectroscopy proved that the contents of different kinds of impurities amount to a few ppm and do not differ significantly in the glasses investigated. Using X-ray absorption spectroscopy the same result was obtained for the contamination of the glass melt with platinum of the melting crucibles. The water contents are very small in the glasses A and D, as revealed by infrared absorption spectroscopy. These measurements were performed with polished glass platelets of 1.5 mm
A. Hoell et aI. / Journal of Non-Crystalline Solids 208 (1996) 294-302
296
thickness. The detection limit for OH- was 10 ppm. Since no OH- signal was registered for either sampies A and D, their hydroxyl content should be lower than a few ppm. The glass transition temperatures, Tg, the densities, d, and the viscosities, r/, of the as-prepared glasses are listed in Table 1. The Tg'S were determined by means of differential thermal analysis with an uncertainty of ___5 K. The measurements of the sample densities were carried out in toluol at 20°C by means of the Archimedes principle. The viscosities were measured at (600 + 1)°C with the fiber elongation method. The measurements were performed with rods of dimension 1 × 1 × 51 mm 3. The binodal temperature, which characterizes the boundary between the metastable and the homogeneous region in the quasi binary phase diagram for a given glass composition, was determined with the opalescence method [15]. Glass strips, about 10 cm long, were heat treated beneath the binodal temperature for a time long enough to ensure the development of an appreciable phase separation throughout the samples. Afterwards the samples were placed in a furnace with a defined, nearly linear thermal gradient. After air quenching to room temperature the glass bars appeared to be homogeneous at the high temperature end and remained phase separated in the temperature region of the miscibility gap below the binodal temperature. The latter is indicated by a strong light scattering of the low temperature side of the glass bar. The binodal temperature corresponds to the position on the bar where the region of maximum opalescence is demarcated from the clear homogeneous glass. Calibrating the temperature field within the furnace this position can be assigned to the corresponding critical temperature with an uncertainty of + 10 K. In order to produce specimens for the SAXS
investigations, the blocks of the glasses A and D were cut into samples of about 25 mm × 10 mm × 3 mm. The as-prepared samples were heat treated isothermally at 600°C for different times. The heat treated samples did not show any crystallization effect. The samples were ground to a final thickness of 0.1 mm and polished using cerium oxide. The surface roughness obtained corresponds to optical quality in order to avoid surface scattering of X-rays. 2.2. SAXS measurements
The X-ray source was a sealed tube with copper anode operating at power settings of 34 kV and 40 mA. Monochromatic Cu K~ radiation was obtained using a Ni filter in combination with a proportional counter with pulse height analyzer. The scattering intensities, l(s), were measured between Smin 0.09 nm -1 and Smax = 3 . 0 nm -1 with a commercial Kratky camera using infinitely long slit collimation conditions [21]. The quantity s denotes the modulus of the scattering vector, s = 41r s i n ( 0 / 2 ) / h , where h is the X-ray wavelength and 0 is the scattering angle. The scattering curve recorded under these conditions contains the complete information about all scattering particles in the size range between 1 nm and 40 nm. All the scattering curves were measured in several runs. The statistical error was less than 3% for any measuring point. After correcting for background scattering the recorded scattering intensities were converted into electron units by means of a calibrated polyethylene (Lupolen) standard sample [21]. From the scattering curves, I(s), the following structural parameters are calculated: the average particle radius, R~, the particle number density, Nv, the internal surface per unit volume, Sv, and the mean square of electron density fluctuations, ((A p)2 ). "------
Table 1 Chemical and physical data of the glasses A and D Glass
A D Batch
Composition (mol%) Na20
CaO
SiO 2
CI
12.13 12.20 13.00
11.02 11.48 11.00
76.38 76.32 76.00
0.47 -
Tg (°C)
Density, d ( g / c m 3)
Viscosity at 600°C, r/(Pa s)
569 570
2.458 2.472
4.0 X l0 II 4.0 × 1011
A. Hoell et al. / Journal of Non-CrystallineSolids 208 (1996)294-302 The mean square of electron density fluctuations, ((A p)2 ), is determined from the first moment of the scattering curve, I(s), by the expression [20] 1
( ( A p ) 2 ) - 2~Aa fo s l ( s ) ds,
(1)
where the quantity a is the distance between the specimen and the detector slit. The scattering intensities recorded between Smin and Smax are extrapolated to zero angle using the Guinier law [20] and to infinity using the Porod law [20]. The Porod law for infinitely long slit collimation is I(s) = k/s 3
(2)
for large values of s. The constant k is connected with the internal surface per unit volume, S,., via
~: = ( ~ / 2 ) A a ( Ap)2S,,.
(3)
For a two-phase system the mean square of electron density fluctuations can be calculated as [20] ( ( A p ) 2 > = w(1 - w ) ( A p ) 2,
(4)
where w and (1 - w) denote the volume fractions of the droplet phase and the matrix phase, respectively, and
A p = UA[(dz/M)droptet- (dz/M)matrix]
R,
foRaN(R)dR
3 foS[(s) ds - f o R 3 N ( R ) d g "
size distribution. It has been demonstrated elsewhere [24] that R t is suitable to test certain power laws predicted by different theories for particle growth during a phase separation process. The particle number density, N,., has been calculated using Eqs. (3) and (6):
St N,
47rR~ "
(7)
3. R e s u l t s
The results of the SAXS investigations indicate considerable differences in the course of the phase separation process in the glasses A and D. This feature is demonstrated in Fig. 1 for the scattering curves of the specimens as-prepared and heat treated at 600°C for 5.5 h. While the untreated sample of glass A already shows an advanced stage of phase separation indicated by an intense small angle scattering effect, no effect due to phase separation is observed in the flat SAXS curve of the untreated sample of glass D. With progressing heat treatment the SAXS effect develops and intensifies. In Fig. 2, the evolution of the structural parame-
(5)
is the difference between the electron densities of the two phases. NA is Avogadro's number, d is the density, z and M are the number of electrons and the molar weight per unit of composition of the two phases. In the present case of the phase separated glass the droplets are enriched with silica at the expense of the silica concentration in the matrix glass. Because of the nearly spherical shape of the droplets the average radius, R 1, which can be extracted from the SAXS data with high accuracy [23], is an appropriate structural parameter for characterizing the particle size. RI can be calculated from the intensities, I(s), according to [20,24]
4 f o I ( s ) ds
297
(6)
As can be seen from Eq. (6), R 1 represents the mean radius of the particle-volume distribution of the droplet phase [23]. N(R) is the corresponding
1027
A (5.5 h) 1026
D (5.5 h) A(0h)
1025 D (0 h)
1024 10-2
~
i
IiA,,,I
. . . .
,,,,I
10o s=4~sin(0/2)/?~ 10-1
,
,
,,,,,,
10 ] Into q]
Fig. 1. SAXS curves of the specimensof glass A (with Cl- ) and glass D (without C1- ) as-prepared and heat treated at 600°C for 5.5 h.
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A. Hoell et al. / Journal of Non-Crystalline Solids 208 (1996) 294-302
lO
20
101 W"
10 0 5.10 5
0
r-------m
~ g
10 5
i
r
i
ilrlll
,
. . . . . . .
till
r
i
1017
T
I
i
i
,~tl
i
,
,
,,,i,I
10
1
D
I
0
100 t [h]
Fig. 3. Mean square of electron density fluctuations, ((Ap)2), and corresponding volume fraction, w, of the nucleus phase in dependence on the time, t, of the isothermal heat treatment. The relative precisions were estimated to be 3% in ((A p)2 ), which do not exceed the extension of the points. The solid lines are drawn as guides for the eye. The value of win,x (dashed line) is calculated from Eq. (8).
.o.
Z
1016
1015
.................. 10
1
100 t [h]
Fig. 2. Average particle radius, R 1, internal surface per unit volume, Sv, and particle number density, N~., of glass A (with C1- ) and glass D (without CI- ) in dependence on the time, t, of the isothermal heat treatment at 600°C. According to [23] the relative precisions were estimated to be 3% in R 1, 8% in S,,, and 14% in N~,. In the case of R], the relative precisions do not exceed the extension of the points.
dashed line in Fig. 3 indicates the theoretical maxim u m value of ( ( A p ) 2) for the case o f complete phase separation, which was calculated using Eqs. (4) and (5) taking into account the metastable i m m i s cibility boundaries experimentally determined b y H a m m e l [15] for the quasi b i n a r y system SiO 2 [ ( 1 3 / 2 4 ) N a 2 0 - ( 1 1 / 2 4 ) C A O ] ( m o l % ) shown in Fig. 4. Here, the m a x i m u m v o l u m e fraction, Wmax , obtain-
1200 o~
ters R 1, S v, and Nv during phase separation at 600°C is shown for glass A containing chloride and glass D without any chloride. O w i n g to fining with 1 wt% NaC1 an appreciable effect on the phase separation behaviour of the glass investigated has been observed. As demonstrated in Fig. 2 a considerable increase of both the particle size and the growth velocity is obtained due to the chloride addition into the melt of glass A. Furthermore, the interface area, S~., and the particle n u m b e r density, N,,, in glass A are larger and they decrease more rapidly than in glass D. Fig. 3 shows the evolution of the m e a n square of electron density fluctuations, ( ( A p ) 2 ) , determined from the SAXS data according to Eq. (1). The
1000 800
o
600 t 400 I I
60
70
f
80
90
100
Xsio= [tool%] Fig. 4. Miscibility gap in the soda lime silica glass according to the quasibinary tie line (data given by Hammel [15]), SiO2[(13/24)NazO-(ll/24)CaO] (mol%). For T=600°C the silica concentrations of the homogeneous glass, Xg = 0.76, of the nuclei, x, = 0.99, and of the matrix, x m = 0.73, are marked. The dashed line indicates the shift of x m and x, caused by a rise of the miscibility gap by 45 K up to higher temperatures.
A. Hoell et al. / Journal of Non-Crystalline Solids 208 (1996) 294-302
able for the glass without any impurity can be calculated for a given treatment temperature from the equilibrium mole fractions of the silica enriched nuclei, x,, and of the matrix, x m, according to the known miscibility gap and the composition of the initial homogeneous glass, Xg, (Xg - x m ) d g M n
(8)
299
proceeds with ripening. Investigating the time dependence of the structural parameters during heat treatment, SAXS offers the opportunity to distinguish between different stages of the phase separation process [3,22,26]. For the chloride containing glass A the following power laws are deduced from Fig. 2:
N, OCt - l ' 1 8 ;
g 1 ~ t0"37;
M and d are the molar weights and the mass densities [25] of the phases denoted by the corresponding indices. Using the experimental data for the miscibility gap given by Hammel [15] the following values are obtained for T = 600°C: Xg = 0.76, x m = 0.73, and x n = 0.99 which are marked in Fig. 4. The density of the initial homogeneous glass was measured to be dg = 2.48 g cm -3. For the purpose of calculating an upper limit for Wmax and for the mean square of electron density fluctuations, ((A p)2), it was assumed that the nuclei consist of pure silica. The corresponding mass density is d n = 2.2 g cm -3. With these data and the Eqs. (4), (5), and (8) the theoretical values of Wmax = 0.13 and ((A p)2 )calc = 7.2 × 1044 cm -6 are obtained and marked by the dashed line in Fig. 3. This figure demonstrates that for long treatment times the experimental ((A p)2) values of the pure glass D should fit asymptotically the calculated value of 7.2 × 1044 cm -6 whereas for glass A, that contains chloride, the equilibrium value of ( ( A p ) 2) is considerably larger and can be estimated to amount to at least 8.9 × 1044 cm -6. Different binodal temperatures were determined for the glasses A and D. The binodal temperature of glass A, Thin,A, amounts to (770 + 10)°C whereas the binodal temperature of glass D, Thin,D, was obtained to be (705 ___10)°C.
If these results are compared with the theoretical predictions of the classical theory of Lifshitz and Slyozov [27] and Wagner [28] (LSW theory) for the asymptotic stage of diffusion-limited ripening,
4. Discussion
4.1. Kinetics The compositions of the glasses A and D lie within the metastable region of the phase diagram. That means that phase separation is going on according to the binodal mechanism, that is characterized by nucleation and growth, which in the late stage
(R)~tl/3;
N,.~t-1;
St, ~ t -0"37.
(9)
Wmax= ( X - Xm)d, Mg
S~,at -1/3,
(10)
it can be concluded that glass A has already reached the asymptotic stage of ripening. The appreciable volume fraction of the nucleus phase, w, which is increasing up to 16.7% in glass A, has no effect on the time exponents in the ripening equations, but has only an influence on the prefactors in the predicted power laws and on the form of the size distribution function [29]. This fact is important because the LSW theory is valid for an infinitely small volume fraction of the droplet phase only. The phase separation kinetics of the pure glass D, in relation to glass A, are considerably slower (Fig. 2), which results in R 1 (xtl/4;
g ~ . ~ t -0"48
S ~ ( x t -0"02 .
(11)
The particle number density, N~,, decreases with progressing treatment time, t, whereas the internal surface, S,., of the scattering system decreases only slightly. It is worth mentioning that in the examined time interval of heat treatment the volume fraction of the precipitating nuclei, w, is less than the saturation value, Wmax (Fig. 3). The power law exponents in (11) for the investigated structural parameters of glass D (Fig. 2) are not typical, neither for the asymptotic stage of diffusion-limited ripening nor for the stage of independent growth. A similar behaviour has already been observed experimentally for a soda lime silica glass investigated by Burnett and Douglas [ 18]. They found a so-called intermediate stage between diffusionlimited growth and the asymptotic stage of
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diffusion-limited ripening. This intermediate stage has a retardation of the increase of the mean particle radius ( R ) during thermal treatment leading to a growth velocity which is temporarily even smaller than in the following stage of asymptotic ripening. Theoretical computations that model the evolution of the nucleus size distribution and the mean particle radius during nucleation, growth, and ripening in viscoelastic media [30] suggest that this phenomenon is typical for phase separating glasses where the relative supersaturation of the initial homogeneous glass in the metastable region of the miscibility gap is relatively large. Apparently the chloride content in glass A acts in such a manner that the phase separation process is accelerated in comparison to glass D which has the same composition ratio concerning the oxide components but does not contain chloride. Therefore, the main result of the kinetic studies is the fact that the chloride containing glass A has already reached the asymptotic stage of ripening whereas for the same treatment conditions glass D is in the intermediate growth stage. Consequently, in case of glass A, the volume fraction and the magnitude of the mean square of electron density fluctuations after the 100 h heat treatment should match with the corresponding saturation values which are valid for the equilibrium case. However, in glass D (without chloride) the maximum phase volume, Win,x (dashed line in Fig. 3), does not precipitate even during a 100 h heat treatment.
4.2. Miscibility gap As demonstrated in Fig. 3, the experimentally obtained magnitude of ( ( A R ) 2) = 8.9 X 1044 c m - 6 for glass A exceeds the theoretical value of 7.2 x 1044 cm -6 calculated from the miscibility gap data of the pure soda lime silica glass by 23%. The influence of the small differences between the compositions of the glasses A and D and, therefore, slight changes of the initial composition in the miscibility gap turns out to correspond to a change o f ( ( A p ) 2) not exceeding +5%. This conclusion was confirmed in Ref. [26]. Obviously, the chloride content in glass A is responsible for the deviation of the final state of phase separation from that expected for the pure
glass. According to Eq. (4) two possible reasons for this effect must be discussed. On the one hand an increased contrast, A p, between the electron densities of the two phases could be caused by the chloride content and on the other hand a higher maximum volume fraction could be developed. Analyzing first a possible change of A/9 one has to consider that the electron density of chloride is very similar to that of the other glass components. The quantitative changes of A p which are obtainable by considering the effect of the chloride content in glass A were estimated to be at least one order of magnitude smaller than necessary in order to explain the increase of ((Ap) 2) observed experimentally. Consequently, the experimental result has to be discussed in terms of an increasing precipitated phase volume, that is larger than w calculated by Eq. (8), owing to the chloride impurity of the glass composition and resulting in a parallel shift of the immiscibility boundary to higher temperatures as illustrated by the dashed binodal curve in Fig. 4. The parallel shift is justified by the results of Kawamoto and Tomozawa [31] who found that the binodal curves of the alkali and earth alkaline silicate glasses can be reduced to one normalized binodal master curve. Furthermore, Kawamoto et al. showed in [32] that the binodal curves of the NazO-B203-SiO 2 system with small amounts of MoO 3 can be represented by curves parallel to the binodal curve of the pure glass. According to Eqs. (4) and (8) the experimentally observed increase of ((Ap) z) by 23% caused by the chloride addition in glass A corresponds to a gap shift of 45 K up to higher temperatures. This result is qualitatively confirmed by the direct measurements of the binodal temperatures of glass A and glass D using a gradient furnace. The thermodynamic driving force, A/z, of a phase separation process depends on the differences in the position of the miscibility gap and the silica concentration of the initial homogeneous glass, Xg, at the temperature of heat treatment, T. Consequently, the increase of the binodal temperature, Tbin, in glass A leads to an increase of A/x which, in a first approximation, is proportional to the undercooling, Tbin -- T [15,26]. Accordingly, the parallel shift of the immiscibility boundary in the quasi binary phase diagram should result in modified kinetic parameters for the whole phase separation process.
A. Hoell et al. / Journal of Non-Crystalline Solids 208 (1996) 294-302
The rate of homogeneous nucleation, given by [33] Inud=loexp(-[AG*
+ AGo]/kT )
/nucl'
is
(12)
with the activation energy for the formation of a critical spherical nucleus 16 cr 3 AG* = - - ~ r 3 A~ 2 "
(13)
301
nucleation and growth in glass A is stimulated by the increased A/x according to Eqs. (13) and (14), whereas an influence of the remaining chloride on the diffusion coefficient, D*, can be to a great extent excluded. Additionally, the higher thermodynamic driving force in glass A results in higher nucleation rates and growth rates (Eqs. (12)-(14)). Consequently, at the end of the stage of independent growth glass A contains more and larger nuclei than glass D (Fig. 2).
The power law for diffusion-limited growth reads [33] (R) 2 =
2D*A/zc~
5. Conclusions t,
(14)
ckT
where k is the Boltzmann constant, T the absolute temperature, AG o the activation energy of diffusion, or the interface tension, D* the effective diffusion coefficient, c= the concentration of the diffusing species in the undisturbed matrix at infinite distance from the separated particle, and c the concentration of the diffusing species in the nucleus. From Eqs. (12) and (13) it can be derived that a rising thermodynamic driving force increases the nucleation rate provided that the activation energy of diffusion, A G o , remains constant. Taking into account the proportionality between the diffusion coefficient, D*, and the reciprocal viscosity of the glass [1], the invariance of A G o can be tested by means of viscosity measurements. For the glasses A and D the same viscosity, 7/= 4.0 × 1011 Pa s, was obtained at 600°C. That means that D* and A G o should have approximately the same values in both glasses. The measured value for matches that determined by Hammel et al. [16] for a 13Na20-11CaO-76SiO 2 glass. These findings correspond to the experiences made by different authors, to which reference is made in Ref. [34], that for soda lime silica glasses in the composition range investigated the viscosity depends only slightly on composition. The identical values of the glass viscosities evidently prove that the differences observed for the phase separation kinetics of the glasses A and D are mainly due to a changed thermodynamic driving force, A/x. The chloride that remained in glass A after the fining is the reason for an increased A/x in comparison to glass D. The observed acceleration of
Small quantities of chloride were proved to remain in the glass melt after fining of soda lime silica glass with sodium chloride. The fined glass A investigated contains 0.47 mol% chloride. This small content of chloride causes changes of kinetics and equilibrium conditions for the phase separation process in the glass in comparison with a pure soda lime silica glass: the phase separation is accelerated and the maximum volume fraction of the droplet silica phase is increased. The effect of chloride is interpreted by a shift of the miscibility gap position to higher temperatures of about 45 K. Viscosity measurements have shown that chloride does not significantly alter the conditions for diffusion in the glass. Thus, the effect of chloride on phase separation is a promotion of the phase separation process caused by a rise of the miscibility gap, which increases the thermodynamic driving force.
Acknowledgements Financial support of the Deutsche Forschungsgemeinschaft (grant Kr 137212-1) is gratefully acknowledged.
References [1] W. Vogel, Glass Chemistry, 2nd Ed. (Springer, Berlin, 1994). [2] D.R. Uhlmann and A.G. Kolbeck, Phys. Chem. Glasses 17 (1976) 146. [3] O.V. Mazurin and E.A. Porai-Koshits, ed., Phase Separation in Glasses (North-Holland, Amsterdam, 1984).
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