The crystal growth and nucleation in glasses in the lithium silicate system have been investigated. Phase separation in ultimately homogenized glasses of the lithium silicate system xLi 2O ·(100 ﹣ x)SiO 2 (where x = 23.4, 26.0, 29.1, and 33.5 mol% Li 2O) has been studied. The glasses of these compositions have been homogenized using the previously established special temperature-time conditions, which make it possible to provide a maximum dehydration and removal of bubbles from the glass melt. The parameters of nucleation and growth of phase separated in homogeneities and homogeneous crystal nucleation have been determined. The absolute values of the stationary nucleation rates I st of lithium disilicate crystals in the 23.4Li 2O ·76.6SiO 2, 26Li 2O ·74SiO 2 and 29.1Li 2O ·70.9SiO 2 glasses with the compositions lying in the metastable phase separation region have been compared with the corresponding rates I st for the glass of the stoichiometric lithium disilicate composition 33.51Li 2O ·66.5SiO 2. It has been found that the crystal growth rate has a tendency toward a monotonic increase with an increase in the temperature, whereas the dependences of the crystal growth rate on the time of low temperature heat treatment exhibit an oscillatory behavior with a monotonic decrease in the absolute value of oscillations. The character of crystallization in glasses with the compositions lying in the phase separation region of the Li 2O-SiO 2 system is compared with that in the glass of the stoichiometric lithium disilicate composition. The conclusion has been made that the phase separation weakly affects the nucleation parameters of the lithium disilicate and has a strong effect on the crystal growth.
The glass forming systems may be regarded as model systems for technical glass ceramics. Many papers on investigation of nucleation in oxide glasses are known.
As can be seen from
Nucleated compound | System | No. references |
---|---|---|
The system Li2O-SiO2 [ | ||
Glasses of the stoichiometric composition | ||
Li2O×2SiO2 | Li2O-SiO2 | [ |
Li2O×SiO2 | Li2O-SiO2 | [ |
Na2O×2CaO×3SiO2 | Na2O-CaO-SiO2 | [ |
2Na2O×CaO×3SiO2 | Na2O-CaO-SiO2 | [ |
BaO×2SiO2 | BaO-SiO2 | [ |
CaO×SiO2 | CaO-SiO2 | [ |
Glasses of the compositions close to stoichiometric ones | ||
Na2O×SiO2 | Na2O-SiO2 | [ |
Na2O×2SiO2 | Na2O-SiO2-Cr2O3 | [ |
BaO×2SiO2 | Na2O-BaO-SiO2 | [ |
Na2O×2CaO×3SiO2 | Na2O-CaO-SiO2 | [ |
Glasses of the non-stoichiometric compositions | ||
Li2O×SiO2 | Li2O-Al2O3-SiO2 | [ |
MgCr2O4 | CaO-MgO-Al2O3-SiO2 | [ |
CaO×MgO×2SiO2 | CaO-MgO-Al2O3-SiO2 | [ |
Na2O×Al2O3×6SiO2 | Na2O-Al2O3-SiO2 | [ |
2SnO×P2O5 | SnO-SnO2-ZnO-P2O5 | [ |
Na2O×ZnO×P2O5 | Na2O-ZnO-P2O5 | [ |
BaO×B2O3×P2O5 | BaO-P2O5-B2O3 | [ |
the crystallization of phase separated glasses cannot be referred to as catalytic. It is necessary to speak about the influence of the phase separation on a further crystallization process rather than about catalysis. Actually, a review of a number of experimental works [
1) The boundary between phases plays a role of an initiating agent.
2) Since, after the phase separation, the composition of a matrix or dispersed phase becomes closer to that of a crystallizing phase, the formation and growth of crystals progress more easily.
3) In one of the phases, the metastable phase precipitates and plays a role of a catalyst for the main crystalline phase.
The purpose of this work is to investigate the crystal growth and nucleation in glasses in the lithium silicate and sodium silicate systems that serve as a basis for the preparation of glass ceramic materials.
The rates of nucleation and growth of crystals and phase separated in homogeneities were studied in glasses of the analyzed compositions xLi2O×(100 − x)SiO2 (where x = 23.4, 26.0, 29.1 and 33.5 mol% Li2O). Since it was established earlier that the crystal nucleation parameters are strongly affected by water [
X-ray powder diffraction analysis was performed on a DRON-2 diffractometer (CuKα radiation; operating voltage 30 kV; current 20 mA; detector rotation rate, 2 deg/min). Differential thermal analysis (DTA) was carried out on a MOM derivatograph (heating rate, 10 K/min; sample weight, 1 g ; galvanometer sensitivity, 1/5; reference sample, Al2O3; platinum crucible). Optical microscopy investigations in transmitted and reflected light
Glass no. | Li2О | SiO2 |
---|---|---|
1 | 23.4 | 76.6 |
2 | 26.0 | 74.0 |
3 | 29.1 | 70.9 |
4 | 33.5 | 66.5 |
were performed on Carl Zeiss Jenaval and Neophot 32 microscopes ( Germany ). Electron microscopy studies were carried out on an EM-125 transmission electron microscope (accelerating voltage, 75 kV). Samples were prepared using the method of celluloid-carbon replicas. The viscosity was measured by the bending method on a Klyuev viscometer. The temperature dependence data on the viscosity was processed by the least squares technique with the conventional computer program for determining the coefficients A, B, and T0 in the approximation according to the Vogel-Fulcher-Tammann equation [
In the 33.5Li2O×66.5SiO2 glass with the composition closest to the stoichiometric composition of the lithium disilicate, excess SiO2 is absent and, therefore, no phase separation is observed. The compositions of the glasses containing 23.4 mol% Li2O (no. 1), 26Li2O (no. 2), and 29.1Li2O (no. 3) lie in the metastable phase separation region of the lithium-silicate system. The samples of glasses no. 1, no. 2, no. 3 and no. 4 were subjected to preliminary heat treatment at temperatures in the range 370˚C - 560˚C for different times. Then, they were held at a development temperature of 600˚C for 10 min. If glass no. 1, no. 2 or 3 is held at the temperature T = 600˚C for a time in the range 0 - 10 h, it remains transparent without visible opalescence in visual examination. If the glass is heat treated for t = 2 h 40 min at temperatures of 400˚C - 560˚C, it remains visually transparent and does not become opalescent. After additional heat treatment at 600˚C for 10 min, the glass acquires bright blue (yellow in transmission) going to milky opalescence. The electron microscopic images of the glasses preliminarily heat treated at temperatures in the range 400˚C - 560˚C and developed at 600˚C for 10 min are displayed in Figures 1(a)-1(c).
We counted the number of traces of particles per unit area in the electron microscopic image NS/S0 and determined the sizes of the maximum traces of particles R. The corresponding results for glass no. 1 are listed in
Т, ˚C | NS, μm−2 | R, Å |
---|---|---|
370 | 90 | 103 |
390 | 108 | 150 |
400 | 112 | 160 |
420 | 120 | 190 |
460 | 134 | 250 |
464 | 131 | 260 |
490 | 118 | 300 |
520 | 94 | 335 |
540 | 74 | 375 |
560 | 60 | 408 |
580 | 48 | 440 |
Here, we assume that the dependence of the number of traces of phase separated droplets per unit area of the sample surface on the heat treatment temperature (at a constant heat treatment time) obeys the same function law as the temperature dependences of the stationary nucleation rate of droplets. Strictly speaking, in order to determine the stationary nucleation rate I(T), initially, it is necessary to determine the number of droplets per unit volume of the glass by the number of traces per unit area with the use of stereological formulas. Then, it is required to determine a variation in the number of droplets per unit volume as a function of the heat treatment time. Thereafter, it is necessary to check that the nucleation rate is stationary and to determine its value. This is not a simple problem, which is complicated by the fact that the droplets in the volume of the glass can have different sizes. The obtained dependence of the number of traces on the temperature for a fixed heat treatment time can differ substantially from the dependence of the stationary nucleation rate of droplets. However, we used a simpler method to approximately evaluate the temperature dependence of the nucleation rate of droplets. Instead of the dependence I(T), we constructed the dependence NS(T) of the number of traces of particles developed for the same specified time of development tdev at Tdev after preliminary heat treatment for the same specified time t at all temperatures T. Since the quantities I and NS are approximately proportional to each other, the curves I(T)/Imax and NS(T)/NSmax are very close to each other and, in order to evaluate the temperature of the maximum nucleation rate, it is sufficient to measure the dependence of the number traces of crystals NS without regarding for their shape and sizes [
where Φ* is the increment of the thermodynamic potential of the glass due to the appearance of a critical (in composition and size) inhomogeneity nucleus, ΦA is the thermodynamic potential of diffusion jumps and transformations during the phase separation, and A is the factor that, like the quantities ΦA and Φ*, depends on the glass composition and temperature more weakly than an exponential function. If the glass composition c0 belongs to the bimodal region at a given temperature, the thermodynamic potential Φ* is relatively high because the inhomogeneities that do not reach some limiting concentration c1 appear to be energetically unfavorable irrespectively of their sizes. When the glass composition belongs to the spinodal region, inhomogeneities with an arbitrarily small difference
where B and c are temperature independent constants and TS is the spinodal temperature. From the condition dI/dT = 0, we find that the maximum of the rate is located at the temperature Tmax satisfying the Equation (3).
, (3)
By assuming that the thermodynamic potential ΦA does not depend on the temperature T, i.e., ΦA = EA, and setting ΦA = 83 kcal/mol (activation energy of oxygen diffusion), at
Under the assumption that the thermodynamic potential ΦA is equal to the activation energy of lithium self diffusion, we obtain
Therefore, the maximum of the ratio NS/NSmax in
indirect because the number of inhomogeneities n' nucleated for the time t'
can differ significantly (and differs) from the number n of large particles growing in the course of the development. The reason is that, after the increase in the temperature to 600˚C, the inhomogeneities with the composition and size that reach critical values at 600˚C are predominantly retained and will develop. The inhomogeneities with the composition that approaches the binodal composition at the corresponding temperature during low temperature heat treatment go beyond the binodal with the increase in temperature to 600˚C and should dissolve in the glass located between inhomogeneities. A number of inhomogeneities can merge together when reaching the critical sizes for a temperature of 600˚C. This dissolution and merging of small inhomogeneities can lead to the formation of larger and stable inhomogeneities with a composition close to the binodal composition at the temperature of 600˚C.
Although the relation between the quantities n and n' is complex, the maximum of the ratio NS/NSmax in
Haller et al. [
In order to compare more rigorously the theory with the experiment, it is necessary to know the spinodal temperature TS. This temperature was evaluated as follows. Quenched glasses no. 1 and no. 2 (without low temperature heat treatments) were held in the gradient furnace at temperatures in the range 300˚C - 800˚C for 2 h. This led to the appearance of the opalescent band along the sample length, which was stable with respect to a further increase in the holding time. This band had a higher intensity in the central part and weaker intensities toward higher and lower temperatures up to its complete visual disappearance. The temperature of the upper edge of the visible opalescence (770˚C for glass no. 1) coincided with the phase separation temperature Tl (770˚C for glass no. 1) determined as the temperature of the disappearance of the visible opalescence after the opalescent part of the rod was displaced in the range of temperatures a priori higher than Tl. The analysis of the electron microscopic images demonstrates that, in going from the high temperature edge of the band to the low temperature edge, the number of particle traces per unit area (or the relative area occupied by particle traces) increases linearly. Initially, this process goes slowly and then accelerates. It is clear because the nucleation rate is minimum in the vicinity of the phase separation temperature Tl and increases as the spinodal temperature TS is approached as a result of the decrease in the thermodynamic potential Φ* in expression (1). When the traces of inhomogeneities begin to cover a larger part of the replica area and to merge together, the increase in the relative area S/S0 occupied by particles at
In order to accelerate the search for the temperature range of crystal nucleation, we used the DTA curves and the temperature dependences of the viscosity, because it was previously shown [
458˚C for glass no. 2, 457˚C for glass no. 3 and 450˚C for glass no. 4. The slope of the temperature dependence of the natural logarithm of the viscosity on the reciprocal of the temperature (
The X-ray powder diffraction data (
It can be seen from the images obtained using the optical microscope in reflected and transmitted light (
no. 3 and in the shape of regular spheres in glass no. 2. The number of crystals increases with an increase in the holding time during preliminary heat treatment.
In order to nucleate lithium disilicate crystals, the samples of initial glasses no. 1 - no. 4 were held at temperatures in the range 420˚C - 520˚C (range of the endothermic effect in the DTA curves). Then, the nucleated crystals were developed (to sizes visible in the optical microscope) at 600˚C for 10 min, i.e., at the temperature corresponding to the ascending branch of the exothermic peak in the DTA curve. The kinetics of homogeneous crystal nucleation was described using the stationary nucleation rate Ist, nonstationary nucleation time τ, and temperature dependences of these quantities.
The stationary nucleation rate Ist was determined from the experimental dependences of the number of crystals n(t) as the slope of the stationary portion of these dependences on the holding time of glasses at each holding temperature. In this work, the rate of stationary homogeneous crystal nucleation was determined by the cross section method (in reflected light) and by direct counting of the number of crystals (in transmitted light). The glass samples were subjected to preliminary heat treatment and rapidly cooled in air to room temperature. Then, the samples were repeatedly heat treated at the development temperature of 600˚C for 10 min. With the aim of examining the samples in reflected light, the surface layer was ground off and the prepared surface was polished. In order to increase the contrast of the boundaries between crystals and glass, the sample surface was etched in a 0.01 M HF solution for 10 s. The number and size of cross sections of crystals per unit surface area were determined from the micrographs with the use of the Neophot optical microscope or directly in the field of vision of the Jenaval microscope.
The number of crystals per unit volume n and the number of cross sections of crystals in the micro section NS are related by the expression n = NS/(SD), where S is the micro section area and D is the diameter of a maximum particle in the micro section. The nonstationary nucleation time τ was evaluated using the induction period tind, which differs from the time τ by the temperature independent factor.
The induction period tind was determined as the intersection point between the continuation of the linear portion of the dependence n(t) and the time axis. Then, we constructed the temperature dependences of all the aforementioned quantities.
The nucleation of the lithium disilicate was observed in all glasses. The dependences of the number of developed crystals n in the nucleation temperature range 400˚C - 520˚C (development at 600˚C for 10 min) were used to determine the stationary nucleation rate and the nonstationary nucleation time at each specific temperature. The induction period tind decreases monotonically with an increase in the temperature. The temperature dependence of the induction period tind(T) allows us to determine the activation energies for nucleation of lithium disilicate crystals Eτ from the formula Eτ = Rdlntind/d(1/T). The corresponding activation energies are equal to 129, 128, 127 and 126 kcal/mol for glasses no. 1 - no. 4, respectively. The temperature dependences of the stationary nucleation rate of lithium disilicate crystals Ist(T) for the glass with the stoichiometric lithium disilicate composition and glasses with displaced compositions are shown in
where Tmelt is the melting temperature and the enthalpy HA, activation entropy SA, and the nucleation barrier correspond to T = Tmax. It can be seen from
relative displacements of structural units that are necessary for crystal nucleation and occur through the breaking and switching of chemical bonds become sufficiently fast with an increase in the temperature beginning from the glass transition point Tg. The presence of the viscosity determined by the free activation energy Φη = ΦA leads to a shift in the maximum of the stationary nucleation rate Ist(T) according to relationship (6) toward higher temperatures (Tmax = Tg = 2/3Tmelt) as compared to its position Tmax = 1/3Tmelt that would be observed at H = 0.
Before proceeding to the discussion of the results of the experiment on the determination of the growth rate of lithium disilicate crystals in glasses no. 1 - no. 3, let us consider the main notions used when determining the crystal growth. The crystal growth occurs as a result of accidental attachments and detachments of structural units from the surface of a supercritical nucleus (the supercritical nucleus is a nucleus with the size exceeding a critical size). The growth of the super-critical nucleus is thermodynamically favorable because this is accompanied by a decrease in the free energy of the system ΔΦ. The linear growth rate of the nucleus is determined by the equation
where l is the thickness of the monomolecular layer growing on the crystal for some time interval, equal to the linear size of the structural unit), and β+ and β− are the probabilities of attachment and detachment of one structural unit from the crystal surface per unit time:
, (8)
Here, τ0 is the time of the order of the period of vibrations of atoms in a chemical bond, whose switching provides the transition of the structural unit to the ordered state of the crystal structure of the nucleus; ΔΦ is the change in the free energy of the system; and a is the linear size of the structural unit. Substitution of these expressions into formula (7) gives
where
Now, we consider the temperature dependence of the stationary growth rate. Under the assumption that
According to expression (10), the growth rate is equal to zero at T → 0 and T → Tmelt. At some temperature (
From the extremum condition
The quantity kTmelt/HA in formula (12) usually does not exceed 1/10. Therefore, the temperature of the maximum growth rate is close to the melting temperature, and it is very difficult to experimentally determine this temperature. This inference is very important for experimenters.
There is a large number of works on the rate of crystal nucleation (see
As can be seen from
Glass no. | Growth temperature, ˚C | ||||||||
---|---|---|---|---|---|---|---|---|---|
570 | 590 | 600 | 640 | 685 | 700 | 718 | 720 | 730 | |
Growth rate U(T), μm/min | |||||||||
1 | 0.045 | 0.125 | 0.25 | 0.68 | 1.25 | - | 1.72 | 1.92 | 2.0 |
2 | 0.036 | 0.125 | 0.31 | 0.44 | 1.12 | - | 1.66 | 1.58 | - |
3 | 0.042 | 0.130 | 0.40 | 0.42 | 1.14 | 1.40 | 1.70 | 1.85 | - |
4 | 0.225 | 0.625 | 1.25 | 3.4 | 6.25 | 8.6 | - | 10.0 | - |
can increase in an uncontrollable manner due to the fact that the time of heat treatment of the sample (1 - 2 min) is comparable to the time of its heating to a specified development temperature and we cannot state with assurance that crystals grew accurately at this temperature. Therefore, although the temperature dependence of the growth rate should have a maximum, we cannot argue that its true position is recorded. It is quite probable that the maximum of the growth rate is located at higher temperatures. The maxima of the nucleation rates are located at lower temperatures than the maxima of the growth rates. Theoretically, they can either be located at a large distance from each other or overlap. Since the accurate position of the maximum of the growth rate is not known, we can only to make the inference that, for the compositions under investigation, the positions of the temperature maxima of the nucleation and growth rates are spaced along the temperature scale by an uncertain value no less than 260˚C (720˚C - 460˚C).
The growth rates of lithium disilicate crystals as a function of the holding time glasses at a constant temperature were investigated at the temperature of 450˚C. The results of the measurements are presented in
It can be seen from this figure that the dependence of the growth rate of lithium disilicate crystals in glasses no. 1, no. 2 and no. 3 on the time of low-temperature heat treatment exhibit an oscillatory behavior with a gradual decay of the oscillation amplitude and a retardation of the growth process. In the framework of the existing theories of nucleation and growth, it was demonstrated for metal alloys [
the crystal growth rate does not depend on the time of low-temperature heat treatment (
The gradual decrease in the crystal growth rate with an increase in the time of holding of the glass at a temperature of 450˚C for glasses no. 1, no. 2 and no. 3 (Figures 8(a)-8(c)) can be explained by the fact that the crystals grow in the glass with the composition lying in the metastable phase separation region. A part of the material is spent for forming the crystal layer, and the growing crystal appears to be surrounded by the region depleted in the building material (the so-called diffusion zone) (
When the droplets come closer to the growing lithium disilicate crystal, they will block its growth, which manifests itself in the dependences (Figures 8(a)-8(c)) as a decrease in the crystal growth rate.
Thus, we have investigated the kinetics of phase separation in three lithium silicate glasses containing 23.4, 26.0 and 29.1 mol% Li2O with the compositions lying in the phase separation region in the Li2O-SiO2 system. It has been found that the temperature dependences of the ratio of the number of traces to the maximum number of traces NS/NSmax and the ratio of the radius of the particle trace to the maximum radius R/Rmax exhibit maxima. The crystal nucleation rate has been studied in glasses with the compositions displaced from the stoichiometric lithium disilicate composition toward an increase in the SiO2 content. The following features have been revealed.
1) The absolute values of the crystal nucleation rate in glasses with the lower content Li2O than that in the glass with the stoichiometric lithium disilicate composition vary insignificantly.
2) The position of the temperature maximum of the nucleation rate for the glass with the lower lithium oxide content is shifted on the temperature scale by 12˚C. This relatively small change in the quantities Ist(Tmax) and Tmax is explained by the fact that the glasses with the displaced composition correspond to the phase separation region in
the Li2O-SiO2 system, so that the lithium disilicate in glasses with displaced composition nucleates in phase separated inhomogeneities with compositions that are displaced toward the lithium disilicate composition and depend weakly on the temperature.
Variations in the crystal growth rate have been investigated as a function of the temperature and time of heat treatment. It has been shown that the temperature dependences of the crystal growth rate tend to an extremum. It has been found that the crystal growth rate depends substantially on the heat treatment time for phase separated glasses and that the growth rate is constant for the glass with the stoichiometric lithium disilicate composition. The dependence of the crystal growth rate on the heat treatment time in phase separated glasses exhibits a damped oscillatory behavior, which is explained by the depletion of the building material for crystals and the necessity of some time for its supply. Therefore, the performed investigations allow us to draw the following conclusion. Unlike the universally accepted opinion, the phase separation not only does not facilitate the conditions for the formation and growth of crystals but, in some cases (see, for example, the dependences of the crystal growth rate on the time of low-temperature heat treatment), even retards the crystal growth process. The data obtained on the nucleation and growth of crystals help to find optimum compositions and temperature-time conditions for the preparation of many glass-ceramic nanomaterials in lithium silicate system, including photostructured materials.
Sycheva, G.A. (2016) Crystal Growth and Nucleation in Glasses in the Lithium Silicate System. Jour- nal of Crystallization Process and Tech- nology, 6, 29-55. http://dx.doi.org/10.4236/jcpt.2016.64004