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New Transparent Glass-Ceramics Based on the Crystallization of “Anti-glass” Spherulites in the Bi2O3−Nb2O5−TeO2 System Anthony Bertrand,†,‡ Julie Carreaud,†,‡ Gael̈ le Delaizir,*,†,‡ Masato Shimoda,§ Jean-René Duclère,†,‡ Maggy Colas,†,‡ Michel Belleil,∥ Julie Cornette,†,‡ Tomokatsu Hayakawa,§ Cécile Genevois,⊥ Emmanuel Veron,⊥ Mathieu Allix,⊥ Sébastien Chenu,†,‡ François Brisset,# and Philippe Thomas†,‡ †

Université Limoges, CNRS, ENSCI, Science des Procédés Céramiques et de Traitements de Surface (SPCTS), UMR 7315, Centre Européen de la Céramique, 87068 Cedex Limoges, France ‡ Labex Sigma-Lim, Université de Limoges, 87032 Limoges, France § Department of Frontier Materials, Nagoya Institute of Technology, Nagoya 466-8555, Japan ∥ Renishaw S.A.S., 77447 Cedex 2 Marne la Vallée, France ⊥ Conditions Extrêmes et Matériaux: Haute Température et Irradiation (CEMHTI), UPR3079 CNRS, 45067 Cedex 2 Orléans, France # Institut de Chimie Moléculaire et des Matériaux d’Orsay (ICMMO), UMR 8182 CNRS, 91405 Cedex Orsay, France S Supporting Information *

ABSTRACT: The crystallization of the Bi0.5Nb0.5Te3O8 (7.14 Bi2O3− 7.14Nb2O5−85.72TeO2) glass composition shows numerous original spherulite-like shape or “droplets” crystalline domains through the bulk. Both SEM-EDS and microprobe measurements demonstrate that both these droplets and the glassy matrix present the same composition. The combination of information obtained from various complementary techniques, electron probe microanalysis, Raman, and SEM-EDS, then leads to the identification of the spherulites as a new Bi0.5Nb0.5Te3O8 “anti-glass” phase. However, the crystallization mechanism is complex since microcracks are observed at the surface of the spherulites in some glass-ceramic materials, suggesting a confined crystal growth. Therefore, the crystallization process appears much different from a homogeneous congruent crystallization. The photoluminescence (PL) properties of the (1 wt %) Er2O3-doped Bi0.5Nb0.5Te3O8 glass-ceramics were investigated during isothermal crystallization of the parent glass at 380 °C. The evolution of the PL signal (4I13/2 → 4I15/2 transition) enables indirect detection of the first steps of crystallization. Moreover, the PL data indicate random nucleation with respect to the location of Er3+ ions, whereas the integrated PL intensity and lifetime values show very comparable evolution trends as a function of the crystallization rate. incorporated into the crystals.4 Glass-ceramics can usually be obtained either through a one-step heat treatment (nucleation and growth at the same time) or via a two-step heat treatment (nucleation followed by the growth of nuclei), leading to different microstructures. In 1983, Burckhardt and Trömel introduced the concept of “anti-glass” that has been reported in tellurite systems.5 An antiglass is defined as a solid material with a clear cationic longrange order but without anionic short-range order. There is therefore order/disorder coexistence in anti-glass materials. Only a few anti-glass structures have been reported. They are

1. INTRODUCTION Tellurite-based glasses present numerous advantages over silica glasses, especially higher transparency in the infrared region up to 5−6 μm, higher refractive index, better solubility of rareearth ions, and promising third-order nonlinear optical properties.1 Due to these interests, TeO2-based glasses may find applications in a variety of optical devices such as optical fibers or optical switches.1,2 The main drawbacks of tellurite glasses remain their poor mechanical properties that can be overcome by a controlled crystallization leading to transparent glass-ceramics.3 Crystallization from glass not only improves the mechanical properties but also, depending on the nature of the crystalline phase, may enhance the nonlinear optical properties as well as the luminescence properties, due to the periodic arrangement of the atoms around the rare-earth ions © XXXX American Chemical Society

Received: July 23, 2015 Revised: August 28, 2015

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DOI: 10.1021/acs.cgd.5b01048 Cryst. Growth Des. XXXX, XXX, XXX−XXX

Crystal Growth & Design

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based on the CaF2 fluorite structure with Te4+ and metal ions statistically distributed in the cationic positions and with the anionic crystallographic positions not fully occupied.5 Anti-glass materials are also called “anion glasses” in reference to the disorder on anion sites. Among reported anti-glasses, SrTe5O11,5 Ln2Te6O11,6 and Bi2Te4O117 materials can be cited. In the cubic Bi2Te4O11 phase, each cation position of the fluorite-type structure is occupied by Bi and Te in a 1/3 to 2/3 ratio, while the anion positions are occupied by oxygen with a 11/12 site occupancy. Luminescence properties have been little investigated in anti-glasses.8−10 The Bi0.5Nb0.5Te3O8 crystalline phase was been studied in 1968 by Meunier et al.11 The structure is similar to the cubic TiTe3O8 structure with lattice parameter a = 11.275 Å (space group Ia3̅). In this work, we propose to study glass-ceramic materials based on the volume crystallization of an anti-glass phase from a vitreous TeO2−Nb2O5−Bi2O3 matrix, with or without the addition of rare-earth Er3+ ions. Peculiar spherulite-shaped crystals or “droplets” are observed following the crystallization process.

temperature T), where m depends on the growth mechanisms and crystal shape. The slope of the plot ln(q) = f(ln(−ln(1 − y))) with q = 5, 10, 15, and 20 °C/min leads to the value of nA. The optimum nucleation parameters were determined from monolithic glass samples using a two-step experiment consisting of a first isothermal step (15 min) at different nucleation temperatures (from 340 to 380 °C) followed by a nonisothermal step until complete crystallization. A monolithic sample and high heating rate (>10 °C/ min) are usually preferred for such a study in order to avoid the formation of nuclei during the heating ramp and also to avoid the nucleation process being a function of the surface of the particles. The images of the glass-ceramics were taken using a Nikon Eclipse 50× optical microscope. The nature of the crystalline phases was determined using a D8 Bruker diffractometer, operated in the θ−2θ configuration (2θ range, 10°−60°) and equipped with monochromatic Cu Kα1 radiation. Both RT and variable temperature (VT) X-ray diffraction (XRD) experiments were carried out. Raman spectroscopy characterizations were performed using an In via Reflex Renishaw Raman spectrophotometer. We used the 2D mapping module in confocal mode which gave access to highdefinition maps. In fact, the large number of spectra recorded (108500 for the 2D map; 300 × 300 μm2 with a step size of 1 μm) enables high spatial resolution images to be generated. The maps were recorded under a 532 nm wavelength excitation and at low power to avoid any damage of the sample (less than 1 mW), using a 50× objective. The use of a 2400 grooves/mm grating permitted reaching a resolution of 1.2 cm−1. The chemical composition of the samples was determined using SEM-EDS Cambridge equipment and electron microprobe measurements (EPMA) (microprobe SX 100 CAMECA). An electron backscatter diffraction (EBSD) map was recorded using a TSL/EDAX system mounted on a FEG-SEM (Zeiss SUPRA 55 VP) system in order to localize the crystalline parts of the glass-ceramic materials. Optical transmission measurements were carried out in the 300− 3300 nm range, at normal incidence, using a Varian Cary 5000 spectrophotometer operated in a dual beam configuration. The linear refractive index, n, of the samples was extracted from ellipsometry measurements, performed using a Horiba Jobin-Yvon UVISEL extended-range spectroscopic ellipsometer operated in the 200− 2100 nm range. Two angles of incidence, fixed at 60° and 70°, were employed for the measurements, and the light spot size was 1 mm in diameter. The photoluminescence (PL) properties were measured at RT, using a Fluorolog 3 spectrofluorimeter commercialized by the Horiba-Jobin-Yvon Co. For the recorded PL emission spectra, the data step was fixed to 0.5 nm.

2. EXPERIMENTAL SECTION Glass samples of Bi0.5Nb0.5Te3O8 (7.14Bi2O3−7.14Nb2O5−85.72TeO2 in molar percent) and Er3+:Bi0.5Nb0.5Te3O8 compositions were synthesized by classic melt-quenching technique in air in a platinum crucible. The TeO2 (Alfa Aesar, 99.99%), Bi2O3 (Sigma-Aldrich, 99.9%), Nb2O5 (Sigma-Aldrich, 99.99%) and Er2O3 (Sigma-Aldrich, 99,999%) raw materials were first weighed in predetermined quantities and heated in a furnace up to 850 °C for 30 min. The samples were poured into an 8 mm diameter mold and annealed at the glass transition temperature, Tg, 10 °C, for 1 h before cooling to room temperature (RT). A pure Bi2Te4O11 composition was prepared from TeO2 and Bi2O3 raw materials from the cooling of the melt as described by Lovas et al.,12 for comparison with the crystals present in the glass-ceramics. Differential scanning calorimetry (DSC) experiments were performed using TA Instruments Q1000 equipment (with a heating rate of 10 °C/min) to investigate the thermal properties, the kinetics parameters, and the optimal nucleation time and temperature by the Ray et al. method.13 The first crystallization energy, Ex1, was determined using both the Ozawa (eq 1)14 and Kissinger (eq 2)15 equations by varying the heating rate (5, 10, 15, and 20 °C/min).

⎛ −E ⎞ ln(q) = ⎜ x ⎟ + C ⎝ RTx ⎠

(1)

⎛ q ⎞ ⎛ −E ⎞ ln⎜ 2 ⎟ = ⎜ x ⎟ + C′ ⎝ Tx ⎠ ⎝ RTx ⎠

3. RESULTS AND DISCUSSION 3.1. Thermal Properties of the Parent Glass. Figure 1a shows the DSC thermograms recorded at a 10 °C/min heating rate for the 7.14Bi2O3−7.14Nb2O5−85.72TeO2 (Bi0.5Nb0.5Te3O8) glass composition (data recorded for both powder and bulk forms), with and without the addition of Er2O3. The glass transition temperature, Tg, is about 350 °C. The glass stability represented by the Dietzel criterion, ΔT = Tx − Tg (Tx being the temperature of the onset of crystallization) decreases from 57 to 45 °C when 1 wt % Er2O3 is added to the parent glass. Then, the increase of the specific surface of the sample from bulk to powder causes a shift of the crystallization peaks toward lower temperatures (from Tx1 = 407 °C for bulk sample to Tx1 = 381 °C for the powdered sample), which is the signature of preferential surface crystallization.16 However, as shown later, the crystallization process is more complex, and a competition between bulk and surface crystallization occurs since crystals are also found through the whole volume of the sample, as displayed in Figure 1b. In this glass-ceramic, the

(2)

with q being the heating rate, Ex the crystallization energy, R the ideal gas constant, Tx the onset crystallization temperature, and C and C′ some constants. The slope of the plot ln(q) (respectively ln(q/Tx2) = f(1000/Tx) allows the determination of the crystallization energy. The Avrami exponent, nA, was determined from the equation proposed by Ozawa:14 ⎛ − mEx ⎞ 1 ln(q) = ⎜ ⎟ − [ln(− ln(1 − y))] + C″ n ⎝ nRTx ⎠

(3)

⎡ d[ln( −ln(1 − y))] ⎤ ⎢ ⎥ = −n d(ln q) ⎣ ⎦

(3′)

or

T

with y = (AT/A) (A is the surface area of the crystallization peak between the beginning and the end of the crystallization, and AT is the surface area between the beginning of the crystallization peak and the B



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Figure 1. (a) DSC curves of the Bi0.5Nb0.5Te3O8 (7.14 Bi2O3− 7.14Nb2O5−85.72TeO2) glass composition: powder sample without Er2O3 (1 wt %), bulk sample without Er2O3 (1 wt %), and bulk sample with Er2O3 (1 wt %). (b) Optical microscope image highlighting the surface crystallization of the sample, but also revealing droplets present within the bulk.

Figure 2. (a) Plot of the inverse of the temperature at the maximum of the crystallization peak (Tp) vs nucleation temperature for glass samples nucleated for 15 min. (b) Plot of the inverse of Tp vs nucleation time for glass samples nucleated at 360 °C.

surface crystallization only proceeds on a few tens of micrometers. 3.2. Crystallization Kinetics and Determination of the Avrami Parameter, nA. To have a better understanding of the crystallization mechanisms in the 7.14Bi2O3−7.14Nb2O5− 85.72TeO2 glass composition, Ex1 has been calculated from the Kissinger and Ozawa (see the Supporting Information) eqs 1 and 2. Mean values of 336 ± 3 and 479 ± 5 kJ/mol were respectively found for the bulk and the powder samples, showing clearly that bulk crystallization is favored in comparison with surface crystallization. Then, the Avrami parameter was determined from eq 3 or 3′. The slope of the plot ln(−ln(1 − y)) as a function of ln(q) is equal to −nA. The Avrami exponent is respectively equal to 3, 3.2, and 3 at 410, 415, and 420 °C (temperatures selected within the first crystallization peak). This value is also in agreement with bulk crystallization (3D crystallization and spherulite shape) if we assume that nA = m (m being the growth parameter). 3.3. Nucleation and Growth Parameters. As shown in Figure 2a, the optimum nucleation temperature is 360 °C as obtained from the plot of the inverse of the temperature at the maximum of the crystallization peak (Tp) versus nucleation temperature in the monolithic glass sample.17 Figure 2b shows the variation of (1/Tp) with nucleation time in monolithic samples that were nucleated at 360 °C. The onset of the plateau in the plot indicates an optimum nucleation time which

corresponds to 60 min. Glass-ceramics were then obtained by different heat treatments under air to vary the number of nuclei: one step, 380 °C from 30 min to 13 h; two steps, 360 °C for 60 min (nucleation) and 380 or 390 °C from 10 min to 2 h (growth). 3.4. Characterizations of Glass-Ceramics. 3.4.1. Optical Properties and Microstructure of Glass-Ceramics. Figure 3a displays photographs of the parent glass and its corresponding glass-ceramics heat-treated at 380 °C for different dwell times. The one-step heat treatment instead of the two-step heat treatment was chosen arbitrarily. As partial crystallization and crystal growth progressively occur, light scattering becomes more and more pronounced and glass-ceramics tend to become opaque as represented in Figure 3a. The parent glass shows a good optical transmittance up to 75% in the visible and near-IR regions. When crystallizing, the maximum of the transmission drops to reach a zero value for the most crystallized glassceramics as shown in Figure 3b. Control of partial crystallization remains challenging for optical applications. Indeed, the crystals’ size must remain smaller than the incident wavelength to avoid too much light scattering in order to retain transparency in the glass-ceramics. Another approach consists of matching the refractive indices of C



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Figure 3. (a) Photographs of parent glass and glass-ceramics obtained by heat treatment at 380 °C for 2, 3, 5, 8, 11, and 13 h. (b) Optical transmittance spectra (thickness, 1.5 mm) of the parent glass (a) and glass-ceramics obtained by heat treatment at 380 °C for (b) 30 min, (c) 1 h, (d) 2 h, (e) 3 h, (f) 4 h, (g) 5 h, (h) 6 h, (i) 7 h, (j) 8 h, (k) 9 h, (l) 10 h, (m) 11 h, (n) 12 h, and (o) 13 h. (c) Refractive indices of the parent glass and a glass-ceramic heat-treated at 380 °C for 8 h.

both the crystals and the glassy matrix which also leads to transparent glass-ceramics18 and ceramics.19 As seen in Figure 4, here the droplet size is much larger than the incident wavelength. The Rayleigh−Gans−Debye scattering would then have to be considered. As well, the difference in the refractive indices between the glass and the crystals must be considered in order to fully understand the optical transmission properties (Figure 3c). After 30 min at 380 °C (one-step heat treatment), the spherulite crystals appear in the glassy matrix with a size smaller than 10 μm. Up to 2 h at 380 °C, the size of the spherulites does not really evolve, whereas the density of the spherulites clearly increases. The size of the spherulites then increases beyond that time, reaching dimensions up to 100 μm in diameter when the dwell time increases to 6−8 h. At some point (10 h), microcracks are observed at the surface of the

crystals (Figure 4k), indicating important mechanical constraints at the interface glass/crystals. When increasing the dwell time at 380 °C, cracks propagate and give rise to strong light scattering, leading almost to a zero optical transmission in the visible and near-IR ranges. Finally, as expected from the study of the nucleation and growth parameters, a two-step heat treatment, i.e., a nucleation plateau at 360 °C for 1 h followed by a growth step at 390 °C for few minutes, leads to a subsequently larger number of nuclei (and consequently a higher number of crystals) as observed in Figure 4m,n. 3.4.2. Structure of the Spherulites. In order to identify the nature of the crystalline phase(s) appearing during crystallization, the annealed samples were first characterized by XRD at RT. Figure 5a gathers the XRD patterns collected for the different samples whose photographs are displayed in Figure 3a. D



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Figure 4. Optical images of the parent glass (a) and glass-ceramics obtained by heat treatment at 380 °C (one step) for (b) 30 min, (c) 1 h, (d) 2 h, (e) 3 h, (f) 4 h, (g) 5 h, (h) 6 h, (i) 7 h, and (j) 11 h, (k) magnification of a cracked spherulite crystal corresponding to sample j, and (l) 13 h. Twosteps glass-ceramic samples heat-treated at 360 °C for 1 h and 390 °C for (m) 10 and (n) 30 min.

After a few hours of treatment at 380 °C (2 h is the minimum time required to detect crystallization, i.e., the most intense Bragg reflection at 2θ ∼ 27.5°), a unique phase, isostructural with the β-Bi2Te4O11 anti-glass phase, can be identified.7 Interestingly, the intensity ratios of the Bragg reflection observed for the crystalline phase differ from those reported for β-Bi2Te4O11 by Lovas et al.12 The incorporation of Nb5+ in the structure, with Nb5+ exhibiting a lower scattering factor compared to Bi3+ and Te4+ cations, could be at the origin of this evolution. In fact, running XRD experiments as a function of increasing temperature provides precious information for the identification of the crystalline phase. In addition, since the DSC curve in Figure 1a shows more than a single crystallization peak, it appears also important to follow in situ the transformations endured with increasing temperature. Therefore, in order to determine the sequence of crystallization, XRD experiments were performed at RT, and from 330 °C up to T = 480 °C, starting from the 7.14Bi2O3−7.14Nb2O5−85.72TeO2 (Bi0.5Nb0.5Te3O8) glass powder. Figure 5b displays the corresponding data. The phase isostructural to the anti-glass cubic β-Bi2Te4O11 phase is appearing first at 330 °C, and as the temperature keeps increasing, the latter clearly transforms mainly into Bi0.5Nb0.5Te3O8, with some minor proportion of αTeO2. Such transformation suggests that the intermediate phase is a low-temperature polymorph of Bi0.5Nb0.5Te3O8, or more probably an anti-glass phase with the same composition as previously observed in tellurite systems.5 The refined lattice parameter performed by the Lebail method is found to be a = 5.64 (1) Å, a value extremely close to the reported value of β-Bi2Te4O11 (a = 5.639 Å).20 Due to the disordered nature of such peculiar crystal phase types, especially on anion sites, it is assumed that the lattice can easily accommodate the presence of Nb5+ ions (r Nb5+ = 78 pm (CN = 6)), without noticing any obvious modification of the lattice

parameter (r Bi3+ = 117 pm (CN = 6); r Te4+ = 90 pm (CN = 4)).21 3.4.3. Chemical Composition of the Spherulites. SEM-EDS and EPMA microprobe measurements were carried out in order to access the chemical composition of both the glassy matrix and the crystalline droplets. The SEM-EDS mapping shows no evidence of different chemical compositions between the glassy matrix and the crystalline phase (droplet/spherulite) (Figure 6a), as no color contrast is observed. Single local points along the axis glassy matrix-spherulite give also the same results. The black-framed SEM image of Figure 6a shows a fresh fracture of a spherulite. No dendrite is observed as usually described in spherulite-shaped crystallization.22,23 The crystal has a smooth aspect. However, the spherulites (or crystalline droplets) are crystallized as depicted by the EBSD map (Figure 6b). The EPMA measurements performed on two different spherulites show also that the concentrations of the Te, Nb, and Bi elements (expressed in at. %) do not vary at all along a line profile which goes across a crystalline droplet (Figure 6c). In addition, the atomic contents were found to be in total agreement with the initial glass composition 7.14Bi2O3− 7.14Nb2O5−85.72TeO2 (i.e., 4.2% Nb, 4.2% Bi, 25% Te, and 67% O). At this point, we can therefore conclude that the chemical composition of the spherulite is similar to that of the glassy matrix, at the micrometer scale of the analysis (step size of 5 μm). To go further, Raman spectroscopy data were also collected. Raman spectra were first recorded on the parent glass (Bi0.5Nb0.5Te3O8 composition) and on both pure β-Bi2Te4O11 crystalline and glass-ceramic (heat treatment of 6 h at 380 °C) materials (Figure 7a). One can first observe that the Raman bands are rather large (as large as that for the initial glass) in the case of the glass-ceramic sample. This experimental observation reflects the existence of structural disorder at E



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Figure 5. (a) XRD patterns collected for the samples depicted in Figure 4: (a) parent glass and glass-ceramics samples heat-treated at 380 °C for (b) 2, (c) 3, (d) 5, (e) 8, (f) 11, and (g) 13 h. Diagram of β-Bi2Te4O11 calculated from structural data obtained by Lovas et al.12 is also reported in red. (b) XRD experiments performed as a function of temperature (at RT and from 330 to 480 °C), starting from a glass powder.

Bi0.5Nb0.5Te3O8 anti-glass phase unambiguously contains some Nb, exactly as the parent glass of the same composition. Analyzing in more detail, Figure 7a reveals some differences in the normalized Raman intensity between the Bi0.5Nb0.5Te3O8 anti-glass phase and the Bi0.5Nb0.5Te3O8 parent glass. Using such differences allows the reconstruction of a 2D Raman map (Figure 7c) based, for instance, on the intensity of the stretching Nb−O mode (extracted from the fit of the 108500 spectra) at 882 cm−1. By doing so, the presence of the crystalline droplets can be perfectly evidenced, almost as clearly as in the optical image displayed in Figure 4h. The 2D map is displayed as a gray color image (the variation of gray color is associated with the Nb−O mode intensity) (Figure 7c), where moderate intensity variations are in the majority observed between the glassy matrix and the core of the spherulites/ droplets. Careful observation of the spherulites located at the focus point reveals that the intensity of the Nb−O vibration band varies notably between the glass matrix and the core of the spherulite. Indeed, clearly larger intensities are detected

short (large distribution of Te−O, Bi−O and Nb−O distances) and medium ranges, in the case of all the glass-ceramics analyzed, thus, proving the anti-glass nature of the detected crystal phase.24 Thus, the Bi0.5Nb0.5Te3O8 crystalline phase reported by Meunier et al. may also exist under the form of an anti-glass phase as evidenced by this study.11 The decomposition of the Raman spectra from the sample heated for 6 h at 380 °C (illustrated in Figure 7b on the average Raman spectrum obtained from the 108500 spectra collected to generate the 2D map) provides crucial information. The attribution of the main bands is given in Table 1 based on previous work conducted in our laboratory.24 In the 500−550 to 950 cm−1 frequency range, the spectrum was decomposed into four Gaussian bands for the β-Bi2Te4O11 crystal phase, whereas in the case of either the parent glass or the analyzed glass-ceramic, five Gaussian bands were required to correctly reproduce the experimental data (Figure 7b). The principal difference lies in the fact that the only band which is absent in the case of the β-Bi2Te4O11 crystal phase is the Nb−O stretching vibration mode at 882 cm−1. Hence, the so-called F



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Figure 6. (a) SEM-EDS mapping of a polished glass-ceramic sample heat-treated at 380 °C for 3 h. The black-framed SEM image corresponds to a fresh fracture of the spherulite. (b) EBSD map collected for that glass-ceramic. (c) Microprobe measurements (EPMA) performed on two spherulites in one given glass-ceramic.

within the core of the droplets (appearing in white contrast) in comparison to the glassy matrix. The most striking feature of such a 2D map remains the pronounced intensity drop (evidenced by the dark contrast) at the spherulite/glassy matrix interface. It particularly shows that such intensity variation occurs within a thickness of less than 5 μm. Several parameters could impact on the Raman intensity of the Nb−O stretching vibration mode, especially the oscillator strength of the vibration that could be related to the Nb−O distance. The variation of the latter would then be the signature

of mechanical constraints in the droplet as testified in Figure 4k, suggesting clearly a confined crystal growth.25 Indeed, the crystallization process appears much different from a homogeneous congruent crystallization.19 3.4.4. Luminescence Properties of Er3+-Doped Glass and Glass-Ceramics. PL measurements provide some indirect information concerning the first steps of the crystallization, in addition to the characterization of the optical performance in terms of light emission by Er3+ ions. Indeed, the PL emission spectra depend on the local environment around the rare-earth ions (here Er3+), as it was for instance documented, in the case G



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of Er3+-doped GeO2−TeO2−Nb2O5−K2O glasses and glassceramics,26 or illustrated for 75TeO2−20ZnO−4Na2CO3− 1Er2O3 raw glasses heat-treated at different temperatures (slightly above Tg and at Tc).27 Thus, it should be possible to follow the progressive transformation of tellurite glasses into glass-ceramics, provided that a sufficient proportion of Er3+ ions are located within the crystals formed (here the spherulites); the PL signal should then be modified in consequence. Figure 8 represents the normalized PL emission intensity of the 4I13/2 → 4I15/2 transition, recorded for our 7.14 Bi2O3− 7.14Nb2O5−85.72TeO2 bulk samples, annealed at T = 380 °C for different times (from 0 to 24 h). For clarity reasons, the figure has been separated in two parts. Figure 8a highlights the “early” stages of the crystallization, with heating time running from 20 min to 2 h. The PL emission spectra remain almost identical to that of the parent glass, up to 1 h 15 min of treatment (in fact, tiny changes can already be noticed for that particular time). After 1 h 30 min at 380 °C, it can be observed that the collected PL emission spectrum clearly differs from the previous one, suggesting modification of the Er3+ environment. All this indirect information related to the structural evolution of the bulk samples with the heating time can be directly confirmed by the HT-XRD data. Indeed, the first Bragg peak was unambiguously detected at 2θ ∼ 27.5°, for the sample heated during 1 h 15 min at 380 °C (Figure 8c), thus corroborating that the environment of Er3+ ions starts to change. Thus, if we now compare these data to the previous XRD data collected for undoped 7.14Bi2O3−7.14Nb2O5−85.72TeO2 glass (Figure 5a), it seems that the incorporation of Er3+ ions in the system slightly accelerates the first steps of crystallization. At a first glance, this would suggest that the nucleation process and/or the early stages of the crystallization are dependent on the presence of Er3+ ions. A careful analysis of the PL properties indicates that such a statement is actually incorrect and that some crystalline droplets will randomly contain Er3+ ions, whereas others will not. For longer crystallization times (superior to the time where the first crystals are detected), the comparison of the optical microscopy images collected for both doped and undoped samples systematically indicates a larger density of spherulites for Er3+-doped tellurite samples, which will coalesce more rapidly and thus lead to a faster opacification of the doped samples, finally suggesting some real impact of the rare-earth ions on the crystallization process. Such an impact was also reflected in the reduction of Tx in the case of Er3+-doped glasses (see Figure 1a). For longer treatments (i.e., over 1 h 30 min), the PL data (Figure 8b) can be separated in two parts. Up to 4 h, the intensity of the 4I13/2 → 4I15/2 transition located at 1558 nm continuously increases with respect to that of the main component at 1532 nm (which was used for the normalization

Figure 7. (a) Raman spectra of the parent glass, the droplet in a glassceramic (380 °C, 6 h), and the pure β-Bi2Te4O11 crystalline phase. (b) Mean Raman spectrum obtained from the 108500 spectra recorded for the 2D Raman map of one given droplet. (c) 2D Raman mapping of the glass-ceramic of interest (380 °C, 6 h) based on the fitted intensity of the Nb−O vibration mode.

Table 1. Raman Assignment of the Main Vibrational Bands Evidenced on the Parent Glass (Bi0.5Nb0.5Te3O8), the Droplet in the Glass-Ceramic, and the Pure β-Bi2Te4O11 Crystalline Phase wavenumber (cm−1)

attribution

315 and 396 430 477 651−652 744−759 882

[TeO3]2− bending mode symmetrical bridges Te−O−Te and Te−O−Nb; Nb−O−Nb and Te−O−Bi vibrational mods nonsymmetrical bridges Te−O−Te and Te−O−Nb; Nb−O−Nb and Te−O−Bi vibrational modes Te−O streching mode in molecule-like TeO2 Te−O stretching mode in [TeO3]2− polyhedra Nb−O streching mode in NbO6 octahedra H



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can be considered as fully crystallized. Therefore, despite increasing the heating time, the environment of Er3+ ions remains identical, explaining why the shape of the PL emission spectrum does not vary anymore. A closer observation of the PL emission spectra reveals some real differences in the shape of the PL spectra, when compared to the PL emission spectrum acquired for the reference Er3+doped Bi2Te4O11 crystal phase. Indeed, the intensity of the external components of the 4I13/2 → 4I15/2 transition, located at ∼1500 and ∼1600 nm, neatly differs: this is likely related to the fact that the Bi0.5Nb0.5Te3O8 anti-glass phase actually differs from the anti-glass phase β-Bi2Te4O11. The evolution of both the normalized PL integrated intensity for the 4I13/2 → 4I15/2 transition and the corresponding lifetime values (extracted from the PL intensity decay curves−not shown here) as a function of the crystallization time is also reported (Figure 9). First, the normalized intensity remains steady up to a heating time of 1 h at 380 °C (see in particular the inset of Figure 9a). The PL intensity starts increasing sensibly after 1 h 15 min, which is then in perfect agreement with both the XRD data and the appearance of the Bi0.5Nb0.5Te3O8 anti-glass crystal phase and the tiny modifications evidenced in the shape of the PL emission spectrum.

Figure 8. Normalized PL emission spectra (4I13/2 → 4I15/2 transition) in the 1400−1700 nm range, recorded at RT (λexc = 975 nm), for the bulk samples (glass and glass-ceramics) after the crystallization step at T = 380 °C. The crystallization time ranges from 0 to 24 h: (a) From the starting glass up to 2 h of heating, (b) from 1 h 30 min to 24 h. The PL signal from the Er3+-doped β-Bi2Te4O11 crystalline phase is also provided for comparison. (c) θ−2θ XRD patterns collected at RT for Er3+-doped samples devitrified at T = 380 °C for times ranging from 1 to 2 h. Figure 9. (a) Integrated PL intensity of the 4I13/2 → 4I15/2 transition versus the crystallization time. Insert: Zoom in of the 1-2 h time range to highlight the early stages of crystallization. Three zones labeled I, II, and III can be evidenced. (b) Lifetime values reported as a function of the crystallization time (λexc = 978.5 nm ; λem = 1545.5 nm). The same three zones labeled I, II, and III can be evidenced.

of the spectra). After that time, the intensity of that component slightly decreases and the overall shape of the PL emission spectra stops varying for samples annealed 11 h and more, at 380 °C. This observation is in complete agreement with the Xray data, which support the fact that the corresponding samples I



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Over the time range before 2 h (labeled zone I in Figure 9), it was rather clear that the main spectral modifications occurred between 1 h 15 min and 1 h 45 min (see Figure 8a). Such modifications of the Er3+ environment are then accompanied by the corresponding pronounced variation of the normalized intensity (see insert of Figure 9a). A drastic increase of the normalized intensity can be then observed between 2 and 8 h of heating time, which corresponds to the zone labeled II. Finally, beyond 8 h of crystallization time (zone III), it is noticed that the PL intensity suddenly drops down. The increase in the PL intensity is mainly related to the concomitant increase in the lifetime values (Figure 9b). The comparison of the two sets of data is actually very striking, as the evolution of the lifetime values reproduces pretty well that of the PL integrated intensity. In particular, the observed evolution in zones labeled II and III is almost identical. Only the region labeled I differs in the sense that the lifetime value monotonously increases up to 1 h and then remains steady between 1 and 2 h, whereas the PL intensity is steady up to 1 h and then increases in between 1 and 2 h. Such variations within zone I remain to be properly investigated and explained. In zone II, the increase in the average lifetime value must be correlated to the simple fact that the corresponding tellurite samples are glass-ceramics, with a volume fraction of crystals which will increase with longer crystallization times. Statistically, Er3+ ions will then simply be more and more incorporated within the crystalline spherulites, as the crystallization time increases. Another parameter that should be taken into account to explain the evolution of the lifetime value is the crystalline quality of the crystal phase: a longer heating time should allow the growth of larger crystallites. In zone III, the sudden diminution of the lifetime value could be attributed, very likely, to the fact that the spherulites are breaking apart (see the formation of fractures clearly evidenced in Figure 4j,k in the case of undoped samples, after 11 h of heating). These fractures would then constitute the main cause for such reduction in the lifetime value and, as a consequence, in the PL intensity. The production of fresh surfaces will indeed tend to strongly lower the lifetime value, as surfaces will correspond to rather distorted environments for Er3+ ions. The aggregation of the rare-earth ions (effect well documented in the literature,28 produced by a too long heating treatment, could also constitute a secondary cause for the reduction of the lifetime value.

the crystals. However, the mechanisms related to the formation of such droplets are rather complex. 2D Raman maps revealed that the glass matrix/spherulite interface is somehow peculiar, in the sense that the intensity related to the Nb−O stretching vibration neatly drops in this spatial region. This variation would then be the signature of mechanical constraints in the crystalline droplet, suggesting clearly a confined crystal growth. The PL emission properties of Er3+-doped glasses and glassceramics were studied. In particular, Er3+ ions were employed as a local probe to provide indirect structural information on the first steps of the crystallization processes. Both PL and XRD data were found to be in excellent agreement. The evolution of the integrated PL intensity and lifetime values as a function of the crystallization time was finally monitored for the Er3+doped samples. Both quantities show very comparable trends, and the increase/decrease in the PL intensity is systematically accompanied by a respective increase/decrease of the lifetime value.

4. CONCLUSION The 7.14Bi2O3−7.14Nb2O5−85.72TeO2 (Bi0.5Nb0.5Te3O8) glass composition and its corresponding glass-ceramics analogues were investigated. Kinetics parameters were further extracted from DSC measurements to determine the mechanisms of crystallization that occur mainly as bulk crystallization, even though a few tens of micrometers are affected by surface crystallization. A combination of Raman spectroscopy, XRD, EBSD, and optical microscopy experiments clearly allowed assessing that the new anti-glass Bi0.5Nb0.5Te3O8 phase, analogous to the βBi2Te4O11 reported by Trömel, crystallizes through the glassy matrix as spherulite- or droplet-like shapes. Optical transmittance of glass-ceramics decreases as crystallization proceeds due to the mismatch of refractive indices as well as some microcracks that appear at the glassy matrix/crystalline droplet interface at some point. EDS-SEM and microprobe measurements revealed an identical chemical composition between the glassy matrix and

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ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.cgd.5b01048. Crystallization energies Ex of bulk and powder samples calculated from the (a) Kissinger and (b) Ozawa methods and (c) calculation of the Avrami exponent nA (PDF)

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AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]. Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS We acknowledge Dr. J. Monnier and Dr. E. Leroy from ICMPE for microprobe measurements. We also acknowledge the financial support of the Limousin Regional Council. This work benefited also from financial support from the French state managed by the National Agency of the Research under both the “Future Investments” program referred to as Labex Sigma-Lim (Grant No. ANR-10-LABX-0074-01) and the HOLIGRALE project (Grant No. ANR-13-BS08-0008-01). REFERENCES

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