In Situ High-Temperature Probing of the Local Order of a Silicate

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In Situ High-Temperature Probing of the Local Order of a Silicate Glass and Melt during Structural Relaxation Mohamed Naji,† Domingos De Sousa Meneses,* Guillaume Guimbretière, and Yann Vaills CNRS, CEMHTI UPR3079, University Orléans, F-45071 Orléans, France ABSTRACT: Direct observation of structural relaxation at molecular scale in network glasses near the glass transition and in the melt is very challenging. Distribution of structural units forming the glass is commonly believed to depend uniquely on composition and temperature. In this paper, we evidence the dynamical structural changes of the silicate network upon structural relaxation by monitoring the temperature and time evolution of its vibrational signature. Just above Tg, the silicate network presents an unexpected time dependence of its relaxation process characterized by two subsequent regimes leading to a disordered equilibrium state. When annealed at 925 K, the network evolves rapidly toward silica and sodium silicate crystalline phases. Evolution models to describe the temperature time dependences of the Qn distributions are reported and discussed. A general picture of the hierarchical character of structural relaxation in network glasses is drawn.

1. INTRODUCTION Because of the complex temperature time-dependence of structural relaxation in glasses and melts, a clear picture of the atomic rearrangement during relaxation processes is still missing.1−10 For the prototypical strong network glass-former SiO2, made of interconnected SiO4 tetrahedra, the glassy network exhibits only one type of (Si−O) chemical bond and a simple thermally activated mechanism of Si−O bond breaking and reformation is generally accepted as a structural origin of the relaxation processes.18,19 In more complex amorphous systems such as fragile glass formers, the types of chemical bonds are multiple and this richness in the glass network connectivity induces heterogeneity in the relaxation dynamics.11−17 In these systems the dynamics of a group of molecules in one region can be orders of magnitude faster than that of others in another region only a few nanometers away.15,17 Some authors of the present paper have recently shown that this heterogeneity is also present in a sodium silicate strong glass as (Na2O)0.27(SiO2)0.73 exhibiting a depolymerized silicate network: in situ high-temperature Brillouin spectroscopy measurements have shown that at the intermediate range order, the relaxation dynamics of the silicate network contain two distinct relaxation time scales.18 These experimental observations are consistent with recent calculations19,20 and ex per im ent s 1 8 , 2 1 of structural relaxation o f t he (Na2O)x(SiO2)1−x system observed with different structural probes. In this frame, one misunderstood key-point of structural relaxation in (Na2O)x(SiO2)1−x system is the dynamic equilibrium between Qn species populations (Qn entities having 4-n nonbridging oxygen (NBO)22). Indeed, following the Modified Random Network Model,23,24 the distribution of the alkali ions in the silica matrix is not uniform at nanometric scale but instead can form an interconnected network of pockets, i.e. © 2015 American Chemical Society

percolation pathways. This model, supported by molecular dynamic simulation,25 implies that percolation pathways are mapped out by the connection of NBO, i.e. by the distribution of Qn populations. It has been also shown from hightemperature NMR that above Tg, the time scale of viscous flow corresponds to the exchange frequency of the Qn units.26−29 Therefore, there exist a real need for a quantitative estimate of the dynamic equilibrium of Qn populations at high temperature. Nowadays, technological advances in spectroscopy and control of high-temperature environments allow in situ measurements.18,30−33 Among the available techniques, the vibrational spectroscopies give access to short-range order information on the dynamical structure of the probed materials. The most common ones are th IR absorption (IRabs)30,34,35 and Raman Scattering (RS),18,31,35 and the most exotic are hyperRaman scattering (HRS)35,36 and IR emissivity (IRemiss).37−39 All these techniques can be combined in order to obtain relevant and precise information on the structure of isotropic materials such as glasses and melt.35 Unfortunately, RS and HRS because of a lack of precise knowledge of Raman scattering cross sections are usually limited in a quantitative estimate of Qn distribution. Next, the IRabs is known to be a very interesting technique for the study of glasses structure,34,40 but unfortunately because of instrumental limitation it is impossible to implant IRabs for very high temperature characterization. At high temperature, the softening of the samples led systematically to the deformation of the surface which precludes quantitative measurements.30 Hopefully, IRemiss is a promising technique for the in situ high-temperature Received: December 9, 2014 Revised: March 16, 2015 Published: March 25, 2015 8838

DOI: 10.1021/jp512234k J. Phys. Chem. C 2015, 119, 8838−8848

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The Journal of Physical Chemistry C probing of short-range order in glasses. This technique, allows one to quantify both the complex dielectric function of a material and its vibrational properties. As a consequence, the analysis of the evolution of the dielectric function versus frequency, temperature, and time provides relevant information on the microscopic properties of materials, including their structural reorganization at short distances. Therefore, the combination of a highly stable and precise homemade setup allowing recording of infrared emissivity41 with a robust analytical model for the reproduction of silicate spectra (this model was proven working for glassy38,42,43 and crystalline39 silicates) gives access to physical meaning band parameters. Finally, through a calculation based on IR bands areas and taking account of the IR activity of modes, one is able to obtain accurate estimates of the Qn population distributions. This procedure (IRemiss measurement, robust analysis, and calculation of Qn distribution) is applied for the first time to characterize the dynamic structural relaxation in glass and melt states. We report in situ high-temperature observations of structural relaxation in a sodium silicate (Na2O)0.27(SiO2)0.73 glass. We measured the isothermal time dependences of the spectral emissivity at three temperatures typical of three temperature ranges exhibiting characteristic structural behavior: in the lowest range of the glass transition (i.e., just below Tg), above Tg, and in the melt far above Tg. We provide a temperature−time dependence of Qn species during the structural relaxation process.

electrical device ensured a good stability of temperature during time acquisition (measured to be less than 3 K over 1 h of measurement). The device was equipped as well with a blackbody reference. It was a commercial Pyrox furnace having a cylindrical cavity of LaCrO3 with an aperture diameter of 8 mm and operating up to 1700 K. Both the sample and the blackbody were placed inside an isolated chamber, which was purged with dry air. The composition of the air was stable enough during the acquisition time to prevent pollution of the emissivity spectra by absorptions due to H2O and CO2 molecules. The thermal flux emitted was collected by the FT−IR spectrometer, which was placed face to face with the purged chamber. The sample holder and the blackbody were positioned on a rotatable stage. The measured sample and blackbody signals have to be background corrected; that is, the spectrometer’s own emission has to be taken into account and subtracted from the interferograms. Thus, an ambient signal was also acquired. For stability reasons, the blackbody furnace was maintained at constant temperature (1280 K). In this spectral range, the silicate glass is opaque and its spectral emissivity is obtained indirectly by using the following expression:

2. MATERIALS AND METHODS 2.1. Glass Synthesis. The chosen glass for this study is the same as in ref 18 and was prepared in exactly the same way. Ingot of (Na2O)0.27(SiO2)0.73 glass was prepared by melt of SiO2 (99.9%) and Na2CO3 (99.9%) in a Pt−Rh crucible at about 1900 K for 4 h (to eliminate all gas bubbles from the melt and obtain homogeneity) and then poured into a stainless-steel mold kept at room temperature. The addition of 27% of sodium leads to a softening of the glassy matrix and significantly lowers Tg (TDSC = 730 K). The as-quenched ingot of glass was g cut by a diamond saw into 10 × 5 × 0.5 mm3 samples with parallel faces, and then optically polished. The chemical composition was determined by inductively coupled plasma atomic emission spectrometry (ICP-AES); the molar concentration was within 1% of the nominal composition. 2.2. Emissivity Measurements. The principle and the measurement method of the device used in this study are identical to those reported elsewhere,44 i.e. the electromagnetic power emitted by a heated sample was compared with the one emitted by a blackbody radiator. Spectra were recorded in the spectral range from 400 to 1600 cm−1 by making use of a Bruker Vertex 80v Fourier transform IR spectrometer and a DTGS detector with a resolution of 4 cm−1. The experimental way to determine directional spectral emissivity at a fixed temperature is directly based on the theoretical definition45 (eq 1), as the ratio of the directional spectral intensity emitted by the glass sample L(ω,T,θ) and that of the blackbody L0(ω,T).

(2)

E=

L (ω , T , θ ) L0(ω , T )

FT[Is(Ts) − IRT(TRT)] FT[Ibb(Tbb) − IRT(TRT)]

E(ω , Ts) =

P(ω , Tbb) − P(ω , TRT) E bb(ω) P(ω , Ts) − P(ω , TRT)

where the subscripts bb, s, and RT stand for the blackbody, the sample, and the room temperature, respectively. In equation 2 FT denotes the Fourier transform, Ii is the measured interferogram, P represents the Planck’s law at Tbb = 1280 K and TRT = 295 K. Ebb is the blackbody emissivity and ω is the wavenumber. The sample temperature Ts is calculated using eq 2 at the Christiansen point, where E is equal to one. The Christiansen point of the glass sample, located at ∼1270 cm−1, is the wavenumber for which it emits like a blackbody. The uncertainty related to this temperature measurement method is estimated to be 38 min, we observe the opposite evolution of the Qn distribution characterized by a slow increase of Q3 with a diminution of Q2 and Q4 until it reaches a plateau after 75 min. Interestingly, the final Qn populations after a complete relaxation are not so far from the ones at the beginning of the quenched glass. When temperature is increased to 925 K, the structural evolution of the glass continues toward a more stable configuration. In this configuration the Qn entities undergo considerable changes, marked by a rapid decrease of the Q3 species accompanied by an important growth of Q4 and Q2, 8843

DOI: 10.1021/jp512234k J. Phys. Chem. C 2015, 119, 8838−8848

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The Journal of Physical Chemistry C

Figure 7. Parameters of Gaussian profiles corresponding to vibrational modes ν1−ν7: frequency, full-width at half-maximum (fwhm), amplitude, and area under the Gaussian bands as a function of the annealing time at 925 K.

of the same order of the characteristic time corresponding to an exponential decrease of the Raman band intensity corresponding to Q2 species, while the slow relaxation dynamics occurs at the same time scale at which the decrease of the Si−O−Si bending vibration is observed in silica. Accordingly, the nature of the silicate network relaxation below Tg was suggested to originate from two-step molecular rearrangements. The first one is rapid and involves the unconstrained silicate tetrahedra forming the sodium Na+ diffusion channels. The second is slow and concerns the silicate tetrahedra far from the sodium percolation channels. Back to the results obtained here, the absence of change in the emissivity spectra suggests that there is no measurable variation of the distribution of Qn entities during the first hour. By combining the results of structural relaxation at long distances18 and the one reported here (at molecular scale), it is conspicuous that below Tg the glass network reorganizes its structure in such a way that there is no visible evolution of the Qn populations. The relaxation process may act then without Si−O bond breaking which explains the weak temperature dependence of the relaxation times measured at the macroscopic scale (Figure 8 in ref 18) and reflects the constant temperature evolution of shear relaxation time.55

until the glass reach its equilibrium expressed by a plateau in the amplitudes time evolution of Qn species.

4. DISCUSSION In the present study we have shown for the first time the polar dynamic evolution of the silicate network in the glass transition region, above Tg, and in the melt state. The organization of this section will follow the above outlined annealing temperatures: below Tg, above Tg, and in the melt. Below Tg, an annealing at 696 K during 1 h does not lead to any visible evolution of the infrared emissivity spectra, consequently at the molecular level there was no observable change in the silicate network. This result seems to be in disagreement with what has been observed by other techniques. Indeed, from Brillouin and Raman scattering measurements, probing macroscopic and nanometric scales, respectively, it has been shown for the same glass composition that structural relaxation of the silicate network occurs with two dynamics: fast, i.e. 3 h, and slow, i.e. 20 h.18 Attribution of relaxation routes was based on a comparison of relaxation times obtained from Brillouin and Raman data. The fast relaxation time scale is 8844

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Consequently, when annealing begins, the glass acquires enough thermal energy to start a restructurization by involving the most constrained Q3 tetrahedral units. In the present case, more than 10% of the Q3 population is impacted in the entire relaxation process. To go into more detail, we also reported in Figure 8a a fit of the time dependencies of the Qn populations with the models reported in Table 1. Table 1. Best Fit Values of the Experimental Data at 864 K Using Models 1 (t < tc) and 2 (t > tc) and at 925 K Using Model 3 T = 864 K

Figure 8. Time evolution of Qn(t) entities during the annealing experiments at 864 K (a) and 925 K (b). Open symbols mark the data while solid lines are best fits to models 1−3. Evolution of ln K* versus annealing time, calculated for annealing experiment at 864 K (c) and 925 K (d); open symbols are the data, while dashed lines are guides for the eyes. Vertical dashed line represents the threshold time tc.

model 1 (t < tc) Qn(t) = snt + Qn

model 2 (t >tc)Qn(t) = cn exp(−(t − tc)/τ)+ Qn864 K

s2 = s4 = −s3/ 2= 6.7 × 10−4

c 2 = c4 = −c3/2 = 0.035

Q3 = 0.697 Q4 = 0.224 Q2 = 0.079

τ = 9.5 min

Q3864 K = 0.722 Q4864 K = 0.211 Q2864 K = 0.067

T = 925 K model 3 Qn(t) = ent + Qn925 K e2 = e4 = −e3/2 = 0.006

Q3925 K = 0.725 Q4925 K = 0.210 Q2925 K = 0.065

Before tc, the quasi linear evolutions of the Qn populations are rather well described by a linear fit: Qn(t) = snt + Qn, where sn corresponds to the rate of consumption/production of Qn species and Qn represents their concentration at the beginning of the experiment. After tc, the time dependence is better reproduced by a model including an exponential decay: Qn(t) = cn exp(−(t − tc)/τ) + Qn864 K, where cn corresponds to the amplitude of the exponential contribution, τ is the relaxation time, and Qn864 K is the Qn population at the new equilibrium of the medium after relaxation. As shown by the quality of the fits and the parameter values reported in Table 1, which ensures physical constraints such as conservation of the number of tetrahedra and pure modification role for the sodium cations (ΣnQn = Σn Qn864 K = 1, consumption/formation rate = n(Q3)/2 = (nQ4) = (nQ2), all the relaxation process satisfies the chemical equilibrium reaction corresponding to eq 9. Furthermore, the reaction constant K*, describing the displacement to the right, can be defined as K* = ([Q4] × [Q2]/([Q3]2)), where [Q4], [Q2], and [Q3] are the concentrations of Q4, Q2, and Q3, respectively. Plotting ln K* vs time is shown in Figure 8c. At tc, the shift to the right of the reaction constant of eq 9 corresponding to local restructuration of the silicate network, i.e. local conversions of two constrained Q3 units into one Q2 and one Q4, is enough to launch a complete relaxation of the whole structure toward a new equilibrium state. The relaxation mechanism acting above t c also follows the chemical equilibrium reaction, but this time there is a shift to the left with a consumption rate of one Q2 and one Q4 for a creation rate of two Q3.The above phenomenological models proposed to characterize the whole temperature dependencies introduce a small set of parameters that allow describing rather well the evolution of the Qn populations. The values of these parameters are certainly related to the initial state of the glass and the annealing temperature; it is expected that they represent quantitative indicators of the thermal history behaved by the glass and its relaxation route toward an equilibrium state. The comparison of their values for glasses with the same composition but having different thermal histories or annealed

Figure 9. Time dependencies of the areas of the modes attributed to Qn units. Note the appearance of the supplementary v5* mode (IQ*4) after 25 min indicating the beginning of the crystallization process.

Moreover, the data reported here complete a missing puzzle in the interpretation of the decrease in the Raman signal corresponding to Q2 units. As the Raman intensity is not directly related to the concentration of the probed molecule but Raman cross-section must also be taken into account, we then give an evidence that the exponential decrease is very likely caused by a change in the environment of Q2 units (channels relaxation), inducing a decrease in their Raman cross-section. Above Tg (annealing at 864 K), our results reveal for the first time an unexpected two regimes of molecular reorganization of the silicate network. The model described in Section 2.3 to simulate the dielectric function allows us to obtain a much more comprehensive picture regarding the phenomena of structural relaxation occurring at this temperature. As shown in Figure 8a, the structural relaxation presents a critical time threshold at tc = 38 min separating two trends: before it, there is a consumption of Q3 units at nearly constant rate and a parallel increase of Q4 and Q2 species; after it, the reverse transformation is observed with a relaxation like production of Q3 units and a parallel decrease of Q4 and Q2 species until the glass reaches a completely relaxed state at this temperature. 8845

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CONCLUSION In conclusion, relaxation dynamics in network glasses have long been considered a hallmark of systems near Tg, above Tg, and in the melt. The central results of our study are (i) a direct experimental account of the molecular reorganization and dynamics in the glass transition region and beyond it for (Na2O)0.27(SiO2)0.73; (ii) a description of the temperature− time dependences of the Qn populations. Near Tg, the structural relaxation at long-range order is accompanied by rotational and translational motion of tetrahedral silicate units at short-range order, without Si−O bond breaking. Above Tg, the Si−O bond breaking and reformation controls the relaxation process: 2regimes were identified characterized by a critical time threshold tc. Before tc, the glass acquires enough thermal energy to launch a local reorganization by converting the most constrained Q3 tetrahedral units into one Q2 and one Q4. After tc, a complete relaxation of the whole structure toward a new equilibrium state takes place accompanied by a consumption of one Q2 and one Q4 for a creation rate of two Q3. In the melt, the system lowers its energy and evolves toward a more ordered state. The dynamic is slow at the beginning affording enough time for nucleation process and launching the crystal growth afterward. The network evolves rapidly toward silica and sodium silicate crystalline phases. Finally, the hierarchical network restructurization described here brings out the heterogeneous relaxation process and demonstrates the existence of dynamically freezed regions during the glass formation.

at other temperatures will be a means to improve our knowledge about their significance and verify their representativeness for the description of the relaxation phenomenon and the melt state. In addition, these results provide a better understanding of the heterogeneous relaxation found in these glasses. The presence of the threshold brings out the heterogeneous character of structural relaxation in strong glass formers that is still debated by many authors.14−17,56−59 Furthermore, it demonstrates the existence of dynamically freezed regions during the glass formation58,60 and suggests the presence of microscopic stressed regions inside a nonrelaxed macroscopic structure. After the critical time threshold, enough degrees of freedom were provided by local relaxation and the entire system can relax toward a new equilibrium state imposed by the temperature bath. In the melt (annealing at 925 K), the system evolves toward a more stable state. This trend is reflected by the strong evolution of the population of the Qn ratios and the appearance of an additional absorption band after 25 min of annealing. Before 25 min, the time dependencies of the Qn species are also well fitted by a linear model: Qn (t) = ent + Qn925 K where en corresponds to the rate of consumption/production of Qn n species, and Q925 K represents their concentration at the beginning of the experiment at 925 K. Their parameter values are also compatible with an evolution explained by an important right shift of the chemical equilibrium reaction (eq 9; Figure 8d). The parameters obtained from the fit reflect a network organization very different from the one obtained after the annealing at 864 K. After 9 min at 925 K, the observed band-narrowing and sharpening is a signature of the launch of an important reorganization of the structure that includes first diffusion and after nucleation and crystallization of new phases. During this process, the appearance of an additional vibrational mode v5* having a small fwhm is a direct evidence of the appearance of a Q4 silica phase. Previous studies based on experiments of X-ray diffraction showed indeed that annealing of a soda-lime glass above the temperature Tg transition, leads to the formation of quartz and devitrite phase.61 Even in the model proposed here, it is no more valid in quantifying the Qn distributions in a two phase system (see Figure 9), one could agree that the crystal growth process results in a consumption of Q3 and a production of Q2 and Q4. Furthermore, the immiscibility diagram of Na2O−SiO2 system shows that the inclusion of 27% of Na2O in the SiO2 network induces a strong shift from the decomposition dome.62 No phase separation is observed as a function of temperature for this composition. However, the results presented here suggest the formation of a silica-rich phase and a sodium silicate phase. Probably, a nucleation process is induced at 864 K followed by a crystal growth at 925 K. Therefore, our results call for a better assessment of the immiscibility diagram of Na2O−SiO2 at high temperatures. Relaxation in the melt can be understood within the framework of the potential energy landscape.2,63,64 So, during a relaxation process the melt reduces its energy by exploring different energy basins. If the system has enough energy to explore different basins it reaches the deepest basin corresponding to the most nearly crystalline state. Moreover, crystal nucleation from a melt is kinetically sluggish, thus we identify the slow dynamic (at the beginning of the experience) to a crystal nucleation in the melt.

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AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]. Present Address †

European Commission, Joint Research Centre (JRC), Institute for Transuranium Elements (ITU), Postfach 2340, 76125 Karlsruhe, Germany. Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS This work was supported by the French Ministry of Research and Education. REFERENCES

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