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Local Disorder and Tunable Luminescence in Sr1−x/2Al2−xSixO4 (0.2 ≤ x ≤ 0.5) Transparent Ceramics Alberto J. Fernandez-Carrion,† Kholoud Al Saghir,† Emmanuel Veron,† Ana I. Becerro,‡ Florence Porcher,§ Wolfgang Wisniewski,∥ Guy Matzen,† Franck Fayon,*,† and Mathieu Allix*,† †

CNRS, CEMHTI UPR3079, Universite Orléans, F-45071 Orléans, France Instituto de Ciencia de Materiales de Sevilla (CSIC-US), Avenida Américo Vespucio s/n, Isla de La Cartuja, 41092 Sevilla, Spain § Laboratoire Léon Brillouin CEA-CNRS UMR12, CEA/Saclay, 91191 Gif-sur-Yvette Cedex, France ∥ Otto-Schott-Institute, Jena University, Fraunhoferstr 6, D-07743 Jena, Germany

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S Supporting Information *

ABSTRACT: Eu-doped Sr1−x/2Al2−xSixO4 (x = 0.2, 0.4, and 0.5) transparent ceramics have been synthesized by full and congruent crystallization from glasses prepared by aerodynamic levitation and laser-heating method. Structural refinements from synchrotron and neutron powder diffraction data show that the ceramics adopt a 1 × 1 × 2 superstructure compared to the SrAl2O4 hexagonal polymorph. While the observed superstructure reflections indicate a long-range ordering of the Sr vacancies in the structure, 29Si and 27Al solid-state NMR measurements associated with DFT computations reveal a significant degree of disorder in the fully polymerized tetrahedral network. This is evidenced through the presence of Si−O−Si bonds, as well as Si(OAl)4 units at remote distances of the Sr vacancies and Al(OAl)4 units in the close vicinity of Sr vacancies departing from local charge compensation in the network. The transparent ceramics can be doped by europium to induce light emission arising from the volume under UV excitation. Luminescence measurements then reveal the coexistence of Eu2+ and Eu3+ in the samples, thereby allowing tuning the emission color depending on the excitation wavelength and suggesting possible applications such as solid state lighting.

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INTRODUCTION Transparent inorganic materials exhibit wide-ranging applications for diverse optical devices due to their ability to transmit light.1 Among them, single crystals are usual high performance candidates in optical and electronic domains.2,3 However, shaping flexibility and high doping levels are essential criteria which are hardly met by single crystal technology. In this context, transparent polycrystalline ceramics appear as an alternative class of optical materials which have attracted much attention recently.4,5 Indeed they combine several advantages over single crystals, including wide composition range, high maximum doping level of the host structure, and large shaping flexibility.5 The production of such materials is usually realized through different sintering approaches including spark plasma sintering (SPS),6 hot (iso)-pressure sintering,7,8 and vacuum9 and microwave sintering.10 Optical isotropy (induced by cubic symmetry or negligible birefringence materials) or a crystallite size much smaller than the incident light wavelength, as described by the Rayleigh-Ganz-Debye particle scattering theory,11 is essential for ensuring transparency in ceramics. However, residual porosity after sintering is a significant drawback preventing material transparency12 and technical challenges remain to reproducibly synthesize fully dense (i.e., © 2017 American Chemical Society

transparent) ceramic materials. Current technologies are limited by complex and time-consuming sintering procedures, accessible compositions, the segregation of doping agents at grain boundaries, and the necessary use of nanometer-scale raw materials which induces expensive elaboration processes along with safety issues.4 The complete and congruent crystallization of glass was recently presented as an innovative and promising synthesis approach for transparent ceramics. This process enables to overcome the major drawbacks of both single crystal and powder sintering technologies.13−16 Transparent polycrystalline ceramics can be elaborated from a parent glass which is fully crystallized during a single thermal treatment. This method requires fulfilling two essential conditions which might be challenging depending on material chemistry: (i) synthesis of a parent glass with the same composition as the aimed ceramic (i.e., aiming for a congruent glass crystallization mechanism), and (ii) retention of transparency during the crystallization process (glass and crystalline phases densities must match to avoid cracks and nanometer scale or oriented microstructure Received: July 25, 2017 Published: November 20, 2017 14446

DOI: 10.1021/acs.inorgchem.7b01881 Inorg. Chem. 2017, 56, 14446−14458

Article

Inorganic Chemistry

thickness using an automatic polisher running with silicon carbide papers (grain size down to 2.5 μm). To obtain the final ceramic materials, the Sr1−x/2Al2−xSixO4 (x = 0.2, 0.4, and 0.5) glass disks were heat treated for 6−7 h in an open air atmosphere furnace at 930, 985, and 1010 °C, respectively, to perform full glass crystallization (5 °C/ min heating and cooling rates). The undoped Sr1−x/2Al2−xSi+xO4 (x = 0.2, 0.4, and 0.5) samples were synthesized following the same method in order to perform accurate structural determination. Characterization Techniques. The glass-transition and crystallization temperatures (respectively, Tg and Tc) of the glasses were determined by differential scanning calorimetry (DSC) using a Setaram multi HTC 1600 instrument. The measurements were performed on 200 mg of glass powder placed in platinum crucibles and heated at a rate of 10 °C/min. Electron diffraction was performed on a Philips CM20 transmission electron microscope operating at 200 keV. The sample was first crushed in ethanol, and a drop of the solution with the small crystallites in suspension was deposited onto a carbon-coated copper grid. Laboratory X-ray powder diffraction (XRPD) data were recorded on a Bragg−Brentano D8 Advance Bruker diffractometer (Cu Kα radiation) equipped with a LynxEye XE detector over an angular range of 10° < 2θ < 120°. High-intensity and high-resolution synchrotron powder diffraction (SPD) data were carried out on the 11BM beamline at the Advanced Photon Source, Argonne National Laboratory, US. Data were collected at room temperature on a spinning sample (60 Hz) over the 0.5−60° 2θ range with a 0.001° step using a λ = 0.413978 Å wavelength. Neutron powder diffraction (NPD) data were acquired on the 3T2 diffractometer at the Laboratoire Léon Brillouin (LLB), Saclay, France. High-resolution data were recorded between 4.5° and 121° (2θ) with a 0.05° step using a monochromatic radiation (1.225 Å). Structural elucidation of the Sr1−x/2Al2−xSixO4 (x = 0.2, 0.4, and 0.5) materials was realized using the Rietveld method24 with the JANA200625 and FULLPROF26 softwares. The microstructure of the raw surface of the transparent ceramics was analyzed by scanning electron microscopy (SEM) using a Merlin compact Zeiss Gemini apparatus. In order to perform more detailed microstructure observations, some crystallized samples were investigated by EBSD technique. Previously, ceramic samples were cut perpendicular to the surface and polished with abrasive slurries down to a diamond paste of 1 μm grain size. A final finish of 30 min using colloidal silica (Logitech Syton Typ SF1 (pH = 13.3, grain size 32 nm)) was applied. All SEM-samples were contacted with Ag-paste and coated with a thin layer of carbon at about 10−3 Pa to avoid surface charging in the SEM. These analyses were performed using a Jeol JSM 7001F SEM equipped with an EDAX Trident analyzing system containing a Digiview 3 EBSD-camera. Electron backscatter diffraction (EBSD) was performed using a voltage of 20 kV and a current of about 2.40 nA. The EBSD-scans were captured and evaluated using the software TSL OIM Data Collection 5.31 and TSL OIM Analysis 6.2. Unreliable data points were removed in all data sets by applying a Confidence Index (CI) filter of 0.1 after performing a grain CI standardization. No further clean-ups which actually modify orientations were applied. Pole figures of textures are presented in multiples of a random distribution (MRD). 29 Si and 27Al magic angle spinning (MAS) nuclear magnetic resonance (NMR) experiments were performed on Bruker Avance III spectrometers operating at magnetic fields of 9.4 and 17.6 T, respectively. The 29Si quantitative MAS NMR spectra were recorded at a spinning frequency of 7 kHz with a pulse duration of 1.1 μs (25° flip angle) and a recycle delay of 60 s. The 27Al quantitative MAS NMR spectra were recorded at a spinning frequency of 30 kHz using a pulse duration of 0.4 μs (15° flip angle) and a recycle delay of 0.5 s. The 27Al 2D MQMAS spectra were acquired at a spinning frequency of 30 kHz using the Z-filter MQMAS pulse sequence.27 All spectra were fitted using the DMFit software.28 Quantum chemical calculations with periodic boundary conditions were achieved using the CASTEP code29 which uses a plane-wavebased density functional theory (DFT). Electron correlation effects

are required in the case of birefringent material). Based on this approach, we have recently reported the synthesis of new transparent gallate, aluminate, and aluminosilicate ceramics such as the BaAl4O713,17,18 Sr3Al2O6,14 SrREGa3O7,19 and Sr1+x/2Al2+xSi2−xO8 (0 < x ≤ 0.4) materials.15 The latter strontium aluminosilicate system is a hexagonal crystalline solid solution located on the SiO2 rich-side of the tie line between SrAl2Si2O8 and SrAl2O4 in the SrO−Al2O3−SiO2 ternary diagram (Figure SI1). Such compositions were reported to be synthesizable at a large scale. The detailed characterization of the structure and the microstructure of the obtained materials demonstrated that the transparency of the ceramics results from the absence of porosity, very thin grain boundaries, and tunable structural disorder which minimizes birefringence effects (hexagonal structure).15 In this work we address the elaboration of new transparent aluminosilicate ceramics located on the alumina-rich side of the SrAl2O4−SrAl2Si2O8 tie line. The synthesis of such materials is a promising challenge as the compositions are close to the commercial but opaque SrAl2O4 material famous for its green long-lasting luminescence properties.20 Indeed, the emission properties of transparent ceramics are expected to be much enhanced compared to classic opaque ceramics prepared by solid state reaction (volume effect in contrast with surface emission for opaque ceramics). However, glass formation of the SrAl2O4 composition remains complex given its high melting temperature and the quenching rate required.21 Herein, a limited substitution of Al2O3 by SiO2 in the initial SrAl2O4 composition is performed to assist 1% Eu-doped Sr1−x/2Al2−xSixO4 glass formation (Sr vacancies have to be considered for charge balance). Europium was used as doping agent to demonstrate the capability of the new compositions to work as luminescent materials. Compositions between 0.2 < x ≤ 0.5 are synthesized; more specifically Sr0.9Al1.8Si0.2O4, Sr0.8Al1.6Si0.4O4, and Sr0.75Al1.5Si0.5O4 compositions are studied in this work. We thus report new transparent ceramic phosphors synthesized by full and congruent crystallization of these glasses via a surface crystallization mechanism. The samples crystallize in a new 1 × 1 × 2 supercell compared to the hexagonal SrAl2O4 crystalline phase. Subsequent detailed structural characterization of the undoped ceramic compositions evidence both specific cationic ordering and remaining local disorder in the structure. These new materials exhibit transparency in the visible region as well as tunable luminescence arising from the entire sample volume.

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EXPERIMENTAL SECTION

Synthesis Procedure. Three Eu-doped (1% Eu) Sr1−x/2Al2−xSixO4 (x = 0.2, 0.4, and 0.5) glass compositions were prepared using CO2 laser-heating coupled to aerodynamic levitation technique.22,23 High purity SrCO3 (99.9%, Strem Chemicals), Al2O3 (99.98%, Alfa Aesar), SiO2 (99.99%, Strem Chemicals), and Eu2O3 (99.99%, Sigma-Aldrich) powders were used (Eu was considered to substitute Sr). For each of the three compositions, stoichiometric amounts of each precursor were weighed and mixed together in ethanol using an agate mortar to ensure homogeneity. The powder mixtures were then pressed into pellets and a small portion of the latter was placed on a metallic nozzle, where it was levitated by an argon flow. Two CO2 laser beams, from top and bottom, were used to ensure homogeneous heating and melting. Glass beads were obtained by free-cooling of the high temperature liquid under contactless conditions after shutting off the laser heating. The obtained glass beads were then annealed at 50 °C below their respective glass transition temperature (Tg) for 1 h to remove internal stress and then optically polished to disks of 1.1 mm 14447


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Inorganic Chemistry were described using the PBE generalized gradient approximation and the core−valence interactions were described by ultrasoft pseudopotentials30 (USPP) generated using the on-the-fly (OTF) generator included in CASTEP. An energy cutoff of 600 eV was used for the plane wave basis set expansion and the Brillouin zone was sampled using a Monkhorst−Pack grid spacing of 0.04 Å−1. Computations of the NMR parameters were performed after DFT-PBE geometry optimization of the x = 0.17 and x = 0.5 structural models. The geometry-optimized structures were obtained by minimizing the residual forces on all atoms up to |F|max below 0.08 eV·Å−1, using P1 symmetry and fixing the cell parameters to the experimentally determined values for x = 0.2 and x = 0.5. DFT calculation of the NMR parameters were performed with the CASTEP code using the projector augmented waves31,32 (PAW) and gauge included projector augmented waves33 (GIPAW) algorithms for computing the EFG and NMR chemical shielding tensors, respectively. The 29Si and 27Al isotropic chemical shifts were deduced from the computed isotropic shielding using previously reported correlations which account for systematic errors on calculated values using this method.34 Transmittance measurements were collected from 200 to 800 nm using a Cary 5000 spectrophotometer equipped with a Photomultiplier and PbS photocell for visible and infrared detection, respectively. The excitation and emission spectra were recorded with a Horiba JobinYbon Fluorolog 3 spectrofluorometer operating in the front-face mode. The CIE chromaticity coordinates of the emitted light were calculated from the emission spectra, considering a 2° observer. The average refractive index of the glass and crystalline materials was determined by fitting the reflectivity spectrum measured with a Bruker Vertex 80 V spectrometer in the infrared range (2−200 μm). From this analysis, an accurate estimation (∼0.02 error) of the refractive index in the visible range is obtained from the physical dielectric function model optimized in the infrared range. The method is not sensitive to the scattering effects that can appear at smaller wavelengths. The XPS measurements were performed on powder samples obtained after crushing the beads in an agate mortar. The spectra were recorded with a SPECS phoybos 100 DLD instrument using non monochromated Al Kα line (1486.6 eV) as the X-ray source. The instrument is equipped with a hemispherical electron energy analyzer working in the constant pass energy mode. The pressure in the analysis chamber was maintained below 5.0 × 10−10 mbar. The Eu 3d region was scanned with a pass energy of 30 eV. The spectra were binding energy calibrated with the Sr 3d5/2 peak at 133.5 eV.

leads to a considerable decrease of the melting temperature (∼1900, ∼1750, and ∼1650 °C for x = 0.2, 0.4, and 0.5, respectively) and eases glass formation. It was thus possible to synthesize transparent glass beads of 3−5 mm diameter via free cooling from the melt. Transparent Eu-doped glasses were also successfully prepared via the same aerodynamic levitation process. The absence of residual crystalline phase in the glass samples was checked using laboratory powder X-ray diffraction (Figure SI2). It is expected that scaled, commercial production of larger glass samples could be attained using a conventional induction or electric arc high-temperature melting industrial process. DSC measurements were then performed in order to determine the working temperature range for further full glass crystallization via an appropriate heat treatment. All the thermograms show a single exothermic peak with onsets at 951, 976, and 985 °C for the x = 0.2, 0.4, and 0.5 Eu-doped compositions, respectively (Figure SI3). Several thermal treatments were thus carried out in order to completely crystallize the samples without overheating the samples and avoid the development of cracks which would prevent transparency of the final ceramic material. The determination of appropriate thermal treatments was assisted by both SEM and XRD in order to track the crystallization mechanism and rate. First, as presented in Figure SI4, a surface nucleation and growth process takes place during the Sr1−x/2Al2−xSixO4 glass crystallization process. Figure 1a, XRD data recorded (i) on the immediate bulk surface, (ii) ca. 200 μm below after polishing, and (iii) on a powder sample demonstrate a strong oriented crystallization mechanism. In order to better understand this behavior and to confirm the high degree of crystallinity of the sample (i.e., absence of glass regions), SEM observations were realized. Figure 1b and Figure SI5, respectively, show the raw and polished surface microstructures of a 1% Eu-doped Sr0.75Al1.5Si0.5O4 sample after annealing at 1010 °C for 7 h. The completely crystallized surface shows crystals with grain sizes ranging between ca. 1 and 4 μm. The EBSD-patterns obtained from the immediate sample surface (Figure SI5b), as well as from polished cross sections (Figure SI5c), were reliably indexed using a material file based on the structure described further in this work. Cross sections were prepared in order to analyze the crystal growth into the bulk. The SEM-micrograph of the polished cross section in Figure 1c shows neither topographical nor material contrast to visualize the crystal growth. However, the superimposed inverse pole figure (IPF) and image quality (IQ)map of an EBSD-scan performed on the area shows that the crystals only grew from the respective surfaces into the bulk until the growth fronts collided in the middle of the sample. As only the top and bottom of this sample were polished, the growth fronts do not form perfect angles toward each other due to the curved sides of the sample. The inset shows the framed area in greater detail to visualize that the crystals change their direction of growth without significantly changing their crystal orientation. This means that the local growth velocity is increased where growth fronts meet, similar to the growth front interaction observed during the crystallization of Sr-fresnoite in a glass.36 The {0001}-pole figures (PFs) 1−3 were calculated from data sets representing 1, the top, and 2, the side of the EBSDscan (Figure 1c). PF no. 3 was calculated from a different cut plane parallel to the initial surface but ca. 200 μm below it. PF 1

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RESULTS AND DISCUSSION Synthesis and Microstructural Characterization of EuDoped Sr1−x/2Al2−xSixO4 (0.2 < x ≤ 0.5) Ceramics. Initially, undoped samples were prepared in order to check the viability of the process. It was found that despite its high melting point (1960 °C),35 the SrAl2O4 (x = 0) composition can be easily melted using a CO2 laser heating coupled to an aerodynamic levitation system. However, it appeared impossible to obtain a SrAl2O4 glass by free cooling (cooling rate of roughly 300 °C s−1 for a 3-mm-diameter bead) from the high temperature liquid (∼2100 °C). Direct crystallization of monoclinic SrAl2O4 (S.G. P21, ICSD 160296 - PDF 01−076−7488) systematically occurred. As recently reported by Shinozaki et al.,21 a much faster (1000 °C s−1) cooling rate is required to ensure glass formation of this composition. In order to improve the glass formation ability of SrAl2O4, the composition was modified along the Al2O3/SrO = 1 tie line of the Al2O3−SrO−SiO2 ternary system (Figure SI1) by substituting small amounts of alumina by silica and reducing the Sr content in appropriate proportions to maintain charge balance like in the fully polymerized SrAl2O4 tetrahedral network (Sr1−x/2Al2−xSixO4 compositions). Increasing the amount of silica in these compositions (6.10, 12.75, and 16.31 SiO2 wt %, respectively) 14448


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Inorganic Chemistry

surface normal is basically random. Similar textures were again observed during the growth of Sr-fresnoite.36 These notable similarities with the growth of Sr-fresnoite justify a closer analysis of the crystal orientation developments near the surface. Figure SI6 shows an SEM micrograph obtained from the cross section of a sample. The superimposed orientation and IQ-map of a performed EBSD scan illustrate that the two highlighted orientation domains initially grew with a homogeneous orientation but incorporated deviations of more than 30° after ca. 100 μm of growth. Hence, the homogeneously oriented crystals near the surface fray into orientation domains which was also observed during the growth of Sr-fresnoite36 and raises questions about the mechanism of crystal growth. While the mechanism of viscose fingering has been proposed to cause this microstructure,37 a more detailed analysis and discussion of this topic is beyond the scope of the work presented here and will be provided elsewhere. The presented SEM and EBSD results show that the glass exhibits a sole surface nucleation while bulk nucleation is not observed. The crystals form highly oriented layers where the primary growth direction is perpendicular to the surface, but the crystallographic c-axes of the crystal lattice are tilted from the surface normal by ca. 40 ± 10°, in agreement with the bulk XRD data (Figure 1a). After 7 h at 1010 °C, the 1% Eu-doped Sr0.75Al1.5Si0.5O4 sample is fully crystallized and residual or noncrystallized glass was not detected. Similar crystallization rates were observed for the x = 0.2 and x = 0.4 compositions. In fact, the same surface crystallization was also observed on the silica-rich side of the SrAl2O4−SrAl2Si2O8 tie line, corresponding to the Sr1+x/2Al2+xSi2−xO8 (0 < x ≤ 0.4) compositions.15 Remarkably, full crystalline materials exhibiting high transparency in the visible range were successfully obtained after heat treatment at 930 °C for 6 h for the x = 0.2 composition, and 985 and 1010 °C for 7 h for the x = 0.4 and x = 0.5 compositions (Figure 2). The highest transparency, varying

Figure 1. (a) X-ray diffraction pattern of the Sr0.75Al1.5Si0.5O4 (x = 0.5) ceramic recorded at the immediate surface, ca. 200 μm below the surface and on a powder sample. (b) SEM micrograph of the raw (nonpolished) surface of the same ceramic. (c) SEM-micrograph of a cross section covering more than half of a sample superimposed by the IPF+IQ-map of an EBSD-scan performed on the area. The framed area is highlighted in the inset. The 0001-PFs were calculated from data sets representing 1, the edge to the polished sample (top) and adjoining growth into the bulk; 2, the growth from the side of the sample; and 3, a cut plane parallel to the initial surface but ca. 200 μm below it (not presented in the figure).

Figure 2. Transmittance curves recorded for Eu-doped Sr1−x/2Al2−xSixO4 (x = 0.2, 0.4, and 0.5) polycrystalline ceramics (1.1 mm thickness). The dotted red line corresponds to the theoretical maximum transmittance calculated to be 91% for an average refractive index of 1.55.

shows a ring texture where the c-axes of the crystals are tilted from the surface normal but rotate around it. The same is indicated by PF no. 2, but here the reference interface is the unpolished side of the sample. PF no. 3 confirms this ring texture and enables us to determine that the c-axes of the crystals are tilted from the surface normal (which is also the primary growth direction) by 40 ± 10° and rotation around the

between 58% and 73% in the visible range (91% is the maximum theoretical transmittance value according to the average refractive index value, n = 1.55), is observed for the x = 0.2 samples. As previously detailed in the case of the Sr1+x/2Al2+xSi2−xO8 (0 < x ≤ 0.4) compositions, this transparency is eased by the absence of porosity and the strong orientation of the birefringent crystals.15 Moreover, as detailed in the next section, the presence of structural disorder in the 14449


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Inorganic Chemistry

direction) relatively to the hexagonal unit cell of the high temperature SrAl2O4 polymorph (P63, a = 8.9291 Å, c = 8.4963 Å).39 The absence of condition limiting the general hkl reflection led to assignment of a primitive (P) lattice. The systematic reflection conditions h − h0l: l = 2n and 000l: l = 2n were observed, suggesting P63mc, P6̅2c, or P63/mcm as possible space groups for these new structures. In order to clarify the origin of the observed extra-reflections and to gain deeper insight on the crystallographic structure of the Sr1−x/2Al2−xSixO4 (x = 0.2, 0.4, and 0.5) transparent ceramics, high resolution synchrotron powder diffraction, neutron diffraction, and both 29 Si and 27Al NMR were performed on undoped samples. First, an autoindexation analysis realized using the TREOR software40 confirmed the 1 × 1 × 2 superstructure (a = 8.75 Å and c = 16.82 Å) and the reflection conditions determined by SAED. For each possible space group, ab initio structure determination was performed from SPD data via the charge flipping method41,42 using the SUPERFLIP program43,44 implemented in the Jana software package.25 Rietveld refinements of synchrotron powder diffraction were then carried out and neutron powder diffraction data were used to complete the determination of the oxygen crystallographic sites using Fourier difference maps and to accurately refine the oxygen positions and thermal parameters. The best agreement between experimental and calculated diffraction patterns was clearly obtained from the P6̅2c space group structural model. Moreover, the models elaborated from the P63mc and P63/ mcm space groups led to inconsistent bond distances and thermal parameters. Therefore, further refinements were conducted in the P6̅2c space group. It should be noted that given the similar X-ray atomic scattering factors and neutron scattering lengths, the site occupancies of the Si4+ and Al3+ sites were not refined. As the Si/Al−O bond distances of the different tetrahedra appeared to be rather similar, mixed sites with statistical Si/Al occupancy matching the nominal composition were considered and fixed along the refinement process. Likewise, given that Si4+ and Al3+ cations occupy the same crystallographic positions, their thermal parameters were constrained to the same values. Moreover, the total amount of Sr was constrained to match the nominal composition. The final refined synchrotron and neutron powder diffraction diagrams for the x = 0.4 composition are shown in Figure 4 and the resulting structural parameters are given in Table 1 and Table SI1. The crystal structures of Sr0.9Al1.8Si0.2O4 (x = 0.2) and Sr0.75Al1.5Si0.5O4 (x = 0.5) were also determined from Rietveld refinements performed on synchrotron powder diffraction and laboratory X-ray diffraction powder patterns, respectively. The structural parameters of the x = 0.4 member were used as a starting model with appropriate Si/Al ratios and Sr contents. Full details of the final structural refinements for these two compositions are given in Supporting Information (Table SI2 and Table SI3). It has to be noted that, while all the reflections could be fitted with the same profile shape on the basis of a = 8.75 Å and c = 16.82 Å hexagonal cell with a P6̅2c space group for the x = 0.5 ceramic (Figure SI8), the structural refinement for the x = 0.2 and x = 0.4 members required an anisotropic broadening model, available in the FULLPROF software,26 for the superstructure reflections, as these appear broader than the SrAl2O4 related reflections (Figure 4a and Figure SI9). Clearly the x = 0.2 material exhibits quite broad and weak extra reflections which sharpen up for the x = 0.4 composition and become almost as sharp as the base structure peak for the x =

material also plays an important role to induce optical isotropy in the anisotropic Sr1−x/2Al2−xSixO4 system, as previously demonstrated for the anisotropic Sr1+x/2Al2+xSi2‑xO4 system.15 Structural Characterization. The structure of the samples was first studied by laboratory X-ray diffraction. Figure 3a

Figure 3. (a) Laboratory X-ray powder diffraction patterns collected on the Sr1−x/2Al2−xSixO4 (x = 0.2, 0.4, and 0.5) ceramics showing the presence of extra reflections for x ≠ 0. In blue, diagram of SrAl2O4 (x = 0) calculated from ICSD 160298 structural file.39 Red arrows indicate peaks corresponding to the 1 × 1 × 2 superstructure. (b) [010]* and [11̅0]* electron diffraction patterns of the x = 0.5 polycrystalline material showing the 1 × 1 × 2 superstructure (a = 8.75 Å and c = 16.82 Å) compared to the hexagonal unit cell of SrAl2O4.

shows that the main reflections of powder XRD data recorded on Sr1−x/2Al2−xSixO4 transparent ceramics (x = 0.2, 0.4, and 0.5) can be indexed using a hexagonal cell (a = 8.9260 Å, c = 8.4985 Å) with P63 space group symmetry, similar to the SrAl2O4 high temperature polymorph (ICSD 160298 − PDF 074−8724). A slight shift of the peaks toward higher diffraction angles is observed along with the Si content increase, in agreement with the Al3+ (0.39 Å) and Si4+ (0.26 Å) ionic radii in fourfold coordination.38 Close inspection of Figure 3a reveals that additional weak intensity reflections can be observed in Sr1−x/2Al2−xSixO4 diffractograms (x ≠ 0 compositions, see inset of Figure 3a) with intensities increasing with x. The same behavior can be observed on Eu-doped compositions (Figure SI7); however, in order to perform accurate structural determination, undoped samples will be further considered in this structure section. Moreover, the presence of the extra reflections could further be confirmed by selected area electron diffraction (SAED) on Sr1−x/2Al2−xSixO4 (x ≠ 0) compositions (Figure 3b). The reconstruction of the reciprocal space led to the identification of a hexagonal unit cell with a = 8.75 Å and c = 16.82 Å, corresponding to a 1 × 1 × 2 superstructure (along the c 14450


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locally and is not well established at the long scale, especially for small substitution values. Figure SI10 shows the evolution of the unit cell parameters as a function of the Si content. The linear unit cell parameter evolution in the Sr1−x/2Al2−xSixO4 system confirms the existence of a solid solution from x = 0.2 to x = 0.5. The cell volume decrease is caused by variations in unit cell dimensions and composition since Si4+ is smaller than Al3+.38 The average P6̅2c structure of the Sr1−x/2Al2−xSixO4 (x = 0.2, 0.4, and 0.5) compounds is depicted in Figure 5. The unit cell of this new hexagonal phase contains 12 formula units with five inequivalent Sr sites with multiplicities 2, 2, 4, 2, 2 (relative to 3 Sr sites with multiplicities of 2 in the 1 × 1 × 1 SrAl2O4 structure), two distinct tetrahedral sites with mixed Al3+ and Si4+ occupancies, as well as five inequivalent O sites (one additional site compared to the hexagonal SrAl2O4 structure). Similarly to the hexagonal SrAl2O4, the structure exhibits layers perpendicular to the c-axis built by rings of six vertex sharing MO4 (M = Al, Si) tetrahedra which point up and down (Figure 5a). Sr2+ ions are located in the cavities of the rings. As mentioned above, the Al3+/Si4+ substitution requires the creation of Sr vacancies for charge balance in the fully polymerized tetrahedral network. Powder diffraction evidence here a specific ordering of the strontium vacancies at the level of the Sr4 (2d) and Sr5 (2a) crystallographic positions of the 1 × 1 × 2 superstructure, with a progressive decrease of the Sr4 and Sr5 site occupancies which reach the values 0 and ∼0.5 at x = 0.5, respectively (Figure 5c). In the average structure determined by powder diffraction, the Sr4 site (similarly to Sr1 and Sr2) is surrounded by 6 T2 tetrahedral sites at ∼3.3 Å and 6 T1 sites at ∼4.0 Å, while the Sr5 site (similarly to Sr3) is neighbored by 6 T1 sites at ∼3.3 Å and 6 T2 sites at ∼4.0 Å. Therefore, the energetically most favorable Al/Si configurations correspond to local charge balance, which requires a Al/Si substitution in the vicinity of the Sr vacancy (3.3 Å), and therefore avoids the formation of Si−O−Si bonds. Such a scenario is expected to influence the Al/Si occupancies of the T1 and T2 sites; however, powder diffraction does not reveal any specific Al/Si ordering in the tetrahedral sites. To further investigate the nature of the Al/Si ordering, 29Si and 27Al solidstate NMR experiments have been employed, which enabled us

Figure 4. (a) Experimental (red circles) and fitted (green solid line) Rietveld refinements of (a) synchrotron and (b) neutron powder diffraction data of the Sr0.8Al1.6Si0.4O4 (x = 0.4) ceramic. Enlargements of the diffractograms are embedded.

0.5 material. We assume that the broadening of the extra reflections is related to the fact that the superstructure is not fully established at a large scale. Indeed, the ordering of the Sr vacancies, related to the organization of the AlO4 and SiO4 species which will be discussed in the NMR section, takes place

Table 1. Atomic Coordinates and Occupancies Determined for Sr0.8Al1.6Si0.4O4 (x = 0.4) from Rietveld Refinement of the Synchrotron Powder Diffraction Pattern Collected at Room Temperaturea

a

at

site

x

y

z

Uiso (Å2)

occupancy

Sr1 Sr2 Sr3 Sr4 Sr5 Al1 Si1 Al2 Si2 O1 O2 O3 O4 O5

2c 2b 4f 2d 2a 12i 12i 12i 12i 12i 6g 6h 12i 12i

0.33333 0.00000 0.66667 0.66667 −0.3197(4) −0.3197(4) 0.6777(3) −0.0021(3) −0.0021(3) −0.0806(6) −0.2846(6) 0.0005(9) −0.1205(9) 0.2106(12)

0.66667 0.00000 0.33333 0.33333 0.00000 0.0033(3) 0.0033(3) 0.3497(4) 0.3497(4) 0.4439(8) 0.00000 0.4173(8) 0.1306(9) 0.4273(5)

0.25 0.25 −0.00390(7) 0.25 0.00000 0.09702(4) 0.09702(4) 0.15228(4) 0.15228(4) 0.0971(3) 0.00000 0.25 0.1446(3) 0.1261(2)

0.0141(5) 0.0125(6) 0.0152(5) 0.0127(5) 0.0089(8) 0.0070(8) 0.0070(8) 0.0118(10) 0.0118(10) 0.0166(5)c 0.0166(5)c 0.0166(5)c 0.0166(5)c 0.0166(5)c

1 1 1 0.207(3) 0.593(3) 0.8b 0.2b 0.8b 0.2b 1 1 1 1 1

Space group P6̅2c, a = 8.79993(1) Å, c = 16.9199(1) Å. bFixed parameter. cParameters constrained to the same value. 14451


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Figure 6. (a) 29Si MAS NMR spectra (black line) of the Sr1+x/2Al2−xSixO4 ceramics (with x = 0.2, 0.4, and 0.5 from top to bottom) recorded at 9.4 T with a spinning frequency of 7 kHz. The red lines correspond to the fit with three individual contributions shown below the spectra (brown, blue, and green lines). (b) GIPAWcalculated 29Si isotropic chemical shifts for the selected structural models with x = 0.167 and x = 0.5 used as benchmark of all possible Si environments in the 1 × 1 × 2 superstructure. The brown up-triangles, blue diamonds, and green down-triangles correspond to Q44Al at a remote distance of the Sr vacancies and to Q44Al and Q43Al in the vicinity of a Sr vacancy, respectively.

Table 2. 29Si Average Isotropic Chemical Shifts, Line Widths, and Relative Intensity of the Distinct 29Si Resonances in Crystalline Sr1−x/2Al2−xSixO4 Samples with x = 0.2, 0.4, 0.5

Figure 5. View along the (a) [001] and (b) [100] directions of the Sr1−x/2Al2−xSixO4 crystal structure. Sr4 and Sr5 sites appear in light green to highlight the presence of vacancies. (c) Occupancy of the Sr2+ sites vs the nominal composition.

to describe the different Si and Al local environments in the structure. As shown in Figure 6a, the 29Si MAS NMR spectra of the Sr1−x/2Al2‑xSixO4 ceramics exhibit several partly resolved 29Si resonances and can be reconstructed considering three individual peaks at −80.4, −82.6, and −85.6 ppm. According to the known 29Si chemical shift ranges,45 these resonances can be assigned to different Qnm units (n and m indicate the number of bridging oxygen atoms and the number of Si−O−Al bonds, respectively) involved in the fully polymerized tetrahedral network (n = 4). Indeed, the resonances peaking at −80.4 and −82.6 ppm correspond to Q44Al units while the remaining one at −85.6 ppm is assigned to Q43Al species. These 29Si MAS NMR spectra thus reveal the presence of Si−O−Si bonds in the network, the amount of which increases with the Al/Si substitution, and directly reflect the deviation from the Loewenstein rule46 stating that Al−O−Al and Si−O−Si linkages are energetically unfavorable in aluminosilicate materials. The observation of two distinct Q44Al resonances in the spectra could be related either to the presence of two distinct tetrahedral sites (T1, T2) in the average structure or to the presence of a Sr vacancy in the Si second coordination sphere as required by local charge balance. Table 2 gathers the average 29Si isotropic chemical shifts, full width at half-

composition

δiso (ppm)

x = 0.2

−80.1 −82.5 −85.7 −80.3 −82.6 −85.6 −80.4 −82.7 −85.9

x = 0.4

x = 0.5

fwhm (ppm) 2.7 2.7 2.9 2.5 2.5 2.9 2.5 2.5 2.9

I (%, ±2) 59 38 3 37 50 13 20 65 15

units 4

Q 4Al Q44Al,[Sr]0 Q43Al,[Sr]0 Q44Al Q44Al,[Sr]0 Q43Al,[Sr]0 Q44Al Q44Al,[Sr]0 Q43Al,[Sr]0

maximum (fwhm), and relative intensities obtained from fits of the MAS NMR spectra. To finally assign the two Q44Al peaks, DFT GIPAW computations of the NMR parameters were performed for various structural models used as a benchmark to describe all the possible Si and Al environments in the superstructure. On the basis of powder diffraction refinements (Sr occupancies), structural models were built for a theoretical x = 0.167 composition for which the 1 × 1 × 2 cell contains 1 Sr vacancy in one of the two Sr4 positions and 2 Si atoms in the 24 T sites, and for the x = 0.5 composition for which the 1 × 1 × 2 cell contains 3 Sr vacancies (2 in Sr4 and 1 in Sr5 positions) and 6 tetrahedral sites occupied by Si atoms. For the 14452


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Inorganic Chemistry two compositions, all inequivalent structural arrangements were generated from the 1 × 1 × 2 average structure using the Supercell program47 and several configurations were arbitrarily selected to span the bottom half of the Coulombic energy landscape. For the selected structures, DFT optimization of all atomic positions was performed prior GIPAW DFT computation of the NMR parameters (the DFT-optimized structural models are shown in Figure SI11). The geometry optimization step, which leads to Al−O and Si−O average bond lengths of about 1.75 and 1.63 Å, appears to be crucial for accurate computation of the NMR parameters (chemical shifts and quadrupolar coupling parameters). The 29Si isotropic chemical shifts calculated for all Si environments in the considered models (Figure 6b) enable us to point out three chemical shift ranges, associated with three distinct Si environments, which nicely reproduce the experimental trends (three partly overlapping 29Si resonances). For each type of Si units, the spread of the calculated chemical shift values reflects the influence of Si−O bond length, Si−Sr distances, and Si−O−T bond angle variations. These results confirm the presence of Q43Al species (peak at −85.6 ppm) and provide unambiguous assignment of the two resonances located at −82.6 ppm and −80.4 ppm to Q44Al units in the close vicinity, and at a remote distance of a Sr vacancy, respectively. Therefore, the 29Si MAS NMR spectra not only prove the presence of Si−O−Si bonds in the tetrahedral network but also enable the quantification of the amount of Si Q44Al units departing from a local charge balance in the network. The extent of Al/Si disorder in the Sr1+x/2Al2+xSi2−xO4 structure was also probed using 27Al 1D MAS and 2D MQMAS NMR experiments at very high magnetic field (17.6 T). When increasing the Si content in the sample, significant modifications of the 27Al MAS spectra are observed, reflecting the formation of Al−O−Si bonds in the network (Figure 7a). For the x = 0.2 sample, 27Al MQMAS experiments evidence two resonances with 27Al average isotropic chemical shifts of 81 and 75 ppm. For the x = 0.4 and 0.5 samples, the MQMAS spectra exhibit two additional resonances with 27Al average chemical shifts of ∼71 and 65 ppm (Figure 7b). These four 27Al individual resonances show asymmetric lineshapes, characteristic of a distribution of the quadrupolar interaction. To a first approximation, the 27Al MAS and MQMAS spectra were simulated with individual lineshapes computed according to the Gaussian Isotropic Model (Figure 7c), in which the distribution of the electric field gradient is assumed to correspond to a statistical disorder,48,49 taken into account a Gaussian distribution of the 27Al isotropic chemical shift. The average 27 Al isotropic chemical shifts, average quadrupolar coupling constants and relative intensities obtained from fits of the 1D and 2D spectra are given in Table 3. Comparison with the spectrum corresponding to SrAl2O4 clearly shows that the peak at 81 ppm corresponds to charge-balanced q44Al units which are the building blocks of the fully polymerized aluminate network. Recently reported 27Al chemical shift trends in calcium aluminosilicates show that the substitution of one Al by one Si in the second coordination sphere of Al results in a −3 ppm shift;50 hence, the contributions at about 75, 71, and 65 ppm would be assigned to q43Al, q42Al, and q41Al units. However, GIPAW computations of the 27Al isotropic chemical shift for the Sr1−xAl2−xSixO4 structural models with x = 0.167 and 0.5 reveal a more complex situation (Figure 7d). Indeed, computational results confirm that the substitution of one Al by one Si in the second coordination sphere of Al results in a

Figure 7. (a) From top to bottom: 27Al quantitative MAS NMR spectra (17.6 T) of the low-temperature (monoclinic) SrAl2O4 phase and the Sr1−x/2Al2−xSixO4 ceramics with x = 0.2, 0.4, and 0.5. Asterisks indicate spinning sidebands of the satellite transitions. (b) 27Al 2D MQMAS spectrum of the sample with x = 0.4 (black lines) and the corresponding simulation with four 27Al contributions. The asymmetric line shape of each individual contribution was reconstructed according to the Gaussian Isotropic Model (GIM) (d = 5 case of the Czjzek distribution).44,45(c) 27Al quantitative MAS spectra of the ceramics (black lines) and their best fits (red lines) with the contributions evidenced in MQMAS spectra. The center bands of the satellite transitions (ST) were taken into account in the simulations. (d) GIPAW-calculated 27Al isotropic chemical shifts for the selected structural models with x = 0.167 and x = 0.5 used as benchmark of all possible Al environments in the 1 × 1 × 2 superstructure. The red down-triangles, circles, and up-triangles correspond to charge-balanced q44Al, q43Al, and q42Al units at a remote distance of the Sr vacancies. The black down-triangles, circles, up-triangles, and squares correspond to q44Al, q43Al, q42Al, and q41Al units in the close vicinity of the Sr vacancies (labeled [Sr]0).

Table 3. 27Al AverAge Isotropic Chemical Shifts (δiso), Average Quadrupolar Coupling Constant (CQ) and Relative Intensity (I) of the Observed 27Al Resonances in Crystalline Sr1−x/2Al2−xSixO4 Samples with (x = 0.2, 0.4, 0.5) composition x = 0.2 x = 0.4

x = 0.5

δiso (ppm) 80.9 74.8 81.3 75.6 71.5 64.0 81.0 75.7 71.5 64.0 59.0

CQ (MHz) 3.78 5.75 4.80 5.95 4.70 5.20 4.70 6.25 4.60 5.20 5.20

I (%, ±2) 59 41 32 45 16 7 17 47 24 9 4

units 4

q 4Al q43Al, q43Al[Sr]0, q44Al[Sr]0 q44Al q43Al, q43Al[Sr]0, q44Al[Sr]0 q42Al[Sr]0 q41Al[Sr]0 q44Al q43Al, q43Al[Sr]0, q44Al[Sr]0 q42Al[Sr]0 q41Al[Sr]0 q40Al[Sr]0

−4 ppm shift, but also highlights that the presence of a Sr vacancy in the second coordination sphere of q4nAl units also leads to a similar shift. This effect of a departure from the local 14453


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Figure 8. (a) Excitation and (b) emission spectra of the 1% Eu-doped Sr0.9Al1.8Si0.2O4 (x = 0.2) transparent ceramic recorded at different emission and excitation wavelengths. (c,d) Emission spectra of the 1% Eu-doped Sr1−x/2Al2−xSixO4 ceramics with x = 0.2, x = 0.4, and x = 0.5 recorded under excitation at 254 nm (c) and at 365 nm (d). The insets in figures (b) and (d) are photographs of the ceramics taken under UV illumination at different wavelengths.

charge compensation for Al q4nAl units mirrors the 29Si chemical shift trends of Q44Al, for which the absence of a Sr vacancy in the second Si coordination sphere leads to a calculated shift of +2 ppm (measured shift of +2.3 ppm). Accordingly, the 27Al peak at 81 ppm is assigned to charge-balanced q44Al units, the broader peak at ∼75 ppm is assigned to both charge-balanced q43Al units and q44Al groups in the close vicinity of a Sr vacancy as well as q43Al without local charge compensation, while the two remaining lines at 71 and 65 ppm are assigned to q42Al and q41Al species. At this stage, the lack of resolution of the 27Al high-field MAS and MQMAS spectra makes it difficult to quantify all the q4nAl units in the network and the associated structural defects relative to the fully ordered atomic arrangement of lowest energy. This limitation could be overcome using additional constraints obtained from high-field 27Al−29Si double-resonance spectral edition experiments.50,51 All in one, while diffraction experiments indicate that Sr1−x/2Al2−xSixO4 compounds (x = 0.2, 0.4, 0.5) adopt a 1 × 1 × 2 superstructure with a long-range ordering of Sr vacancies, solid-state NMR shows that some degree of disorder remains within the Al/Si tetrahedral network. The extent of disorder can be quantified from 29Si MAS NMR spectra, which unambiguously reveal the presence of two types of structural defects, associated with departure from the Lowenstein rule (excess of Si−O−Si and Al−O−Al bonds) and deviation from local charge balance in the fully polymerized network, respectively. The 29Si NMR spectra show that the amount of Si−O−Si bonds (which remains much smaller than expected assuming a simple Al/Si binomial distribution) increases with the Si content, while the relative content of Si Q4nAl units departing from local charge compensation decreases with the Si

content (i.e., increasing the number of Sr vacancy in the structure). The degree of disorder in the tetrahedral network of these ceramics obviously depends on the annealing time and temperature of the parent glasses. These results thus reveal different long-range ordering kinetics in the strontium and tetrahedral subnetworks, the latter involving the minimization of the amounts of Si−O−Si and Al−O−Al bonds while satisfying local charge balance of Q4nAl and q4nAl species. Luminescence Properties of the Eu-Doped Sr1−x/2Al2−xSixO4 Transparent Ceramics. In order to evaluate the luminescence properties of the new transparent Eu-doped Sr1−x/2Al2−xSixO4 ceramics, we have recorded excitation and emission spectra for the 1% Eu-doped phosphors. The excitation spectra were recorded by monitoring the characteristic emissions of Eu3+ and Eu2+ at 614 and 475 nm, respectively. The excitation spectrum recorded for the x = 0.2 sample at λem = 614 nm (Figure 8a, red line) displayed a set of features in the 300−400 nm range, which are due to the direct excitation of the electrons from the Eu3+ ground state to higher energy levels of the 4f-manifold, the most intense one appearing at 393 nm. Moreover, it is also possible to observe the tail of an intense band at low wavelength values ascribed to a charge transfer process (O2− → Eu3+).52 The excitation spectra recorded for the x = 0.4 and x = 0.5 ceramics at λem = 614 nm are identical to this one, as observed in Figure S12a. On the other hand, the excitation spectrum recorded for the x = 0.2 ceramic at λem = 475 nm (Figure 8a, blue line) exhibits a very broad band with a maximum at around 335 nm which can be assigned to the 4f7 → 4f65d1 (4f−5d) transition of Eu2+ ions.53 The excitation spectra recorded for the x = 0.4 and x = 0.5 ceramics at λem = 475 nm (Figure S12b) are also similar to 14454


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respectively, resulting in a Eu3+ → Eu2+ reduction mechanism as the negative charges in vacancy defects would be transferred into Eu3+ by thermal stimulation. The last requirement states that (4) the host compound must have an appropriate tetrahedral anion environment (like, for example, BO4, SiO4, SO4, AlO4, or PO4) to favor the Eu3+ → Eu2+ reduction process described above.64 Such a configuration creates a compact 3D network around the Sr 2+/Eu 2+ sites which have been demonstrated to effectively shield the oxidation of Eu2+ ions.59,65 As previously described, the structure of our Sr1−x/2Al2−xSixO4 transparent materials exhibits ring layers built by six vertex sharing MO4 (M = Al, Si) tetrahedra where the Sr2+/Eu2+ sites are placed inside the cavities of the rings, thus preventing the oxidation of Eu2+ ions. In summary, the Sr1−x/2Al2−xSixO4 system meets the four conditions described above, which explains the observation of Eu2+ in addition to Eu3+ excitation bands in the spectrum of Figure 8a. We also stipulate that the local disorder induced by the Si4+ for Al3+ substitution in the Sr1−x/2Al2−xSixO4 ceramics provides favorable conditions for the stabilization of Eu2+ in the structure. According to the excitation spectra, it is possible to selectively excite Eu2+ by selecting a wavelength where no excitation band of Eu3+ is observed, as for λexc = 365 nm. It is also possible to solely excite Eu3+ at λexc = 254 nm, as no excitation band of Eu2+ appears at this energy. Last, an excitation at λexc = 393 nm will simultaneously excite both Eu3+ and Eu2+. Figure 8b shows the emission spectra obtained after excitation of the x = 0.2 Sr1−x/2Al2−xSixO4 ceramic at 365, 254, and 393 nm. The emission spectrum recorded while exciting the sample at 365 nm shows a broad emission band centered at 450 nm and extending from 400 to 570 nm (Figure 8b, blue line), which is attributed to the 4f65d1 → 4f7 transition of Eu2+.52 Upon excitation at 254 nm, the sharp emission bands emerging between 570 and 730 nm are assigned to the wellknown 5D0−7FJ (J = 0, 1, 2, 3, 4) transitions of Eu3+ (Figure 8b, red line). The emission at around 617 nm, due to the electric dipole transition (5D0−7F2), appears stronger than that of the magnetic dipole transition at 586 nm (5D0−7F1), as expected for Eu3+ ions located in noninversion symmetry sites.52 This result is in good agreement with the crystal structure reported in Table 1. Remarkably, the excitation at 393 nm enables the excitation of both Eu2+ and Eu3+ which then induce multiple emissions over the visible range (Figure 8b, green line). The inset of Figure 8b shows the corresponding emission colors obtained from the x = 0.2 1% Eu-doped Sr1−x/2Al2−xSixO4 transparent ceramic excited at 393, 365, and 254 nm. As expected from the corresponding emission spectra, the ceramic emits red or blue/green light after excitation at 254 or 365 nm, respectively. Interestingly, the ratio between the emission bands of Eu2+ and Eu3+ can be tailored as a function of the excitation wavelength, which results in different emission colors. For example, the color coordinates after excitation at 393 nm are x = 0.29 and y = 0.29, which are placed fairly close to the ideal white light region.66 It is important to notice that using a near UV LED (393 nm) with a single phase white phosphor is advantageous, as it is expected to produce a higher value for the color rendering index compared with that obtained with the traditional blue LED and yellow phosphor.67 In summary, the color coordinates of the x = 0.2 1% Eu-doped Sr1−x/2Al2−xSi+xO4 transparent ceramic can be easily tuned from green to red via the white region by simply changing the value of the excitation wavelength.

this one, although slight changes in the positions of the maxima and in the bandwidth can be observed. The photoluminescence (PL) excitation spectra indicate, therefore, that both Eu3+ and Eu2+ coexist in the aluminosilicate ceramic, even though the synthesis was not performed under reducing atmosphere. The Eu3+/Eu2+ ratio was estimated from the XPS spectra of the samples. Figure S13 shows the experimental XPS spectra (Eu 3d region) corresponding to the Eu-doped Sr1−x/2Al2−xSixO4 samples with x = 0.2, x = 0.4, and x = 0.5. The three spectra clearly show both valence states of Eu (Eu3+ and Eu2+) with their corresponding spin−orbit splitting (3d5/2 and 3d3/2) plus a plasmon peak on each region (labeled on the spectra), in good agreement with the literature.54 The Eu 3d5/2 region was used to evaluate the Eu3+/Eu2+ ratio in the samples. This region was fitted to two Gaussian-Lorentzian curves corresponding to the Eu3+ 3d5/2 and Eu2+ 3d5/2 photolines and an additional line corresponding to the plasmonic peak. The individual contributions and the fitted curves are displayed as well in Figure S13. The percentage of Eu3+ and Eu2+ estimated from the fittings are given in Table 4. It can be observed that the Table 4. Percentages of Eu3+ and Eu2+ Estimated from XPS Spectra of the Sr1−x/2Al2−xSixO4 Ceramics composition

% Eu3+

% Eu2+

x = 0.2 x = 0.4 x = 0.5

75 74 71

25 26 29

Eu3+/Eu2+ ratio is around 3 for all three samples, the slight differences among them being likely due to the low signal-tonoise ratio of the spectra, which is in turn a consequence of the low Eu doping level (1%). However, the effect of a local disorder induced by Si4+ for Al3+ substitution cannot be disregarded. The reduction from Eu3+ (use of a Eu2O3 precursor) to Eu2+ in air has been reported for several matrices such as borates,55 phosphates,56 sulfates,57 aluminates,58 silicates,59 aluminosilicates,60−62 and glass materials.63 Pei et al.64 proposed the next four conditions that seem to be necessary for the reduction of Eu3+ to Eu2+ in solid state compounds when prepared in air atmosphere at high temperature: (1) no oxidizing ions should be present in the host, which is the case for the strontium aluminosilicate matrix; (2) the trivalent doping RE3+ ions must substitute the different divalent cations of the host; and consequently (3) the substituted cations must have similar radii to the divalent RE2+ ions. These last two requirements are fulfilled by our system since Eu3+ and Sr2+ ions have similar radii for all coordination numbers.38 The Eu3+ for Sr2+ substitution mechanism has been widely explained by the model of charge compensation expressed by the following equations:65 3Sr 2 + + 2Eu 3 + → V″Sr + 2[EuSr ]. V″Sr →

x V Sr

.

+ 2e

(1) (2)

x

2[EuSr] + 2e → 2[EuSr]

(3) 2+

According to this model, three Sr ions could be replaced by two Eu3+ ions and one VSr″ negative vacancy defect with two negative charges, which will compensate the induced two positive defects of [EuSr]· to maintain the charge balance. VSr″ and [EuSr]· act as donors and acceptors of electrons, 14455


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The luminescence of the two other compositions is very similar to the one shown above for the x = 0.2 sample when excited at 254 nm, as shown in Figure 8c. However, the emission spectra recorded under excitation at 365 nm on the three 1% Eu-doped Sr1−x/2Al2−xSixO4 ceramics (Figure 8d) show a clear blue-shift with increasing x, so that the color of the emission is greenish on the x = 0.2 sample and blue on the x = 0.4 and x = 0.5 samples. In summary, the emission color of the 1% Eu-doped Sr1−x/2Al2−xSixO4 ceramics can be tuned from red to blue/ green by changing the excitation wavelength on any given composition or from green to blue by changing the x value of the Sr1−x/2Al2−xSixO4 system of the ceramic excited at 365 nm.

The manuscript was written through contributions of all authors. All authors have given approval to the final version of the manuscript. Funding

FOCAL ANR-14-CE07−0002, Equipex Planex ANR-11EQPX-36, MAT2014 54852-R. Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS The authors thank Domingos De Sousa Meneses and Sandra Ory from CEMHTI lab for refractive index determination and thermal measurements, respectively. J.P. Espinós, from the Materials Science Institute of Seville, is gratefully acknowledged for help with XPS. The authors also thank the French ANR for its financial support to the FOCAL ANR-14-CE07-0002 and Equipex Planex ANR-11-EQPX-36 projects, the CNRS and CSIC for the PICS n° 07490 and PIC2016FR1 projects, as well as the Spanish Ministry of Economy and Competitiveness for the grant MAT2014 54852-R. A. J. F.-C. gratefully acknowledges an F. P. D. I. grant from Junta de Andalucia.́ Financial support from the TGIR RMN THC FR3050 for conducting the research is gratefully acknowledged. Use of the Advanced Photon Source at Argonne National Laboratory was supported by the U.S. Department of Energy, Office of Science, Office of Basic Energy Sciences, under Contract No. DE-AC0206CH11357. Neutron experiments were performed at the Laboratoire Léon Brillouin in Saclay (France).

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CONCLUSION Full and congruent oriented crystallization from glass was successfully applied to develop new Eu-doped Sr1−x/2Al2−xSixO4 (x = 0.2, 0.4, and 0.5) ceramics which exhibit remarkable transparency in the visible region of the electromagnetic spectrum. The determined crystalline structure (P6̅2c with a ≈ 8.75 Å and c ≈ 16.82 Å) corresponds to a 1 × 1 × 2 superstructure compared to the hexagonal unit cell of the high temperature SrAl2O4 polymorph. It is characterized by five different Sr2+ crystallographic sites, two distinct tetrahedral sites with mixed Al3+ and Si4+ occupancies, and five O sites. The experimental synchrotron and neutron powder diffraction data indicate that the Al/Si substitution promotes the presence of vacancies on the Sr4 and Sr5 positions. 29Si and 27Al solid-state NMR experiments unambiguously reveal the presence of two types of structural defects, associated with departures from the Lowenstein rule (excess of Si−O−Si and Al−O−Al bonds) and from local charge balance in the fully polymerized network. Luminescence spectroscopy data clearly demonstrate the coexistence of both Eu2+ and Eu3+ in the crystal structure which enables tuning of the emission color by varying the excitation wavelength. The simultaneous excitation of Eu2+ and Eu3+ with the adequate wavelength leads to white light emission, which may pave the way to a new generation of optically active transparent ceramics which could be integrated into light-emitting devices.

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(1) Zhuang, Y.; Ueda, J.; Tanabe, S. Multi-color Persistent Luminescence in Transparent Glass Ceramics Containing Spinel Nano-Crystals with Mn2+ Ions. Appl. Phys. Lett. 2014, 105, 191904. (2) Keller, U. Recent Developments in Compact Ultrafast Lasers. Nature 2003, 424, 831−838. (3) Takagi, K.; Fukazawa, T. Cerium-Activated Gd2SiO5 Single Crystal Scintillator. Appl. Phys. Lett. 1983, 42, 43−45. (4) Goldstein, A.; Krell, A. Transparent Ceramics at 50: Progress Made and Further Prospects. J. Am. Ceram. Soc. 2016, 99, 3173−3197. (5) Wang, S. F.; Zhang, J.; Luo, D. W.; Gu, F.; Tang, D. Y.; Dong, Z. L.; Tan, G. E. B.; Que, W. X.; Zhang, T. S.; Li, S.; Kong, L. B. Transparent Ceramics: Processing, Naterials and Applications. Prog. Solid State Chem. 2013, 41, 20−54. (6) Kim, B. N.; Hiraga, K.; Morita, K.; Yoshida, H. Spark Plasma Sintering of Transparent Alumina. Scr. Mater. 2007, 57, 607−610. (7) Podowitz, S. R.; Gaumé, R.; Feigelson, R. S. Effect of Europium Concentration on Densification of Transparent Eu:Y2O3 Scintillator Ceramics Using Hot Pressing. J. Am. Ceram. Soc. 2010, 93, 82−88. (8) Seeley, Z. M.; Dai, Z. R.; Kuntz, J. D.; Cherepy, N. J.; Payne, S. A. Phase Stabilization in Transparent Lu2O3:Eu Ceramics by Lattice Expansion. Opt. Mater. 2012, 35, 74−78. (9) Huang, Y.; Jiang, D.; Zhang, J.; Lin, Q. Fabrication of Transparent Lanthanum-Doped Yttria Ceramics by Combination of Two-Step Sintering and Vacuum Sintering. J. Am. Ceram. Soc. 2009, 92, 2883−2887. (10) Cheng, J.; Agrawal, D.; Zhang, Y.; Roy, R. Microwave Sintering of Transparent Alumina. Mater. Lett. 2002, 56, 587−592. (11) Apetz, R.; van Bruggen, M. P. B. Transparent Alumina: A LightScattering Model. J. Am. Ceram. Soc. 2003, 86, 480−486. (12) Messing, G. L.; Stevenson, A. J. Toward Pore-Free Ceramics. Science 2008, 322, 383−384. (13) Allix, M.; Alahraché, S.; Fayon, S.; Suchomel, M.; Porcher, F.; Cardinal, T.; Matzen, G. Highly Transparent BaAl4O7 Polycrystalline

ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.inorgchem.7b01881. SrO−Al2O3−SiO2 ternary diagram; X-ray diffraction data; DSC thermograms; Optical and SEM microscopy; Evolution of the Sr1−x/2Al2‑xSixO4 cell parameters; DFTgeometry-optimized structural models; Excitation and XPS spectra; Interatomic distances and atomic coordinates obtained from synchrotron powder diffraction data (PDF)

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REFERENCES

AUTHOR INFORMATION

Corresponding Authors

*E-mail: [email protected]. *E-mail: [email protected]. ORCID

Mathieu Allix: 0000-0001-9317-1316 14456


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