Qilin Zou

Cécile Marcelot

Nicolas Ratel-Ramond

|| Xiaodong Yi

Pierre Roblin

Florian Frenzel

Ute Resch-Genger

Ali Eftekhari

Aude Bouchet

Christophe Coudret

Marc Verelst

Xueyuan Chen

|| Robert Mauricot

Clément Roux

Heterogeneous oxysulfide@fluoride core/shell nanocrystals for upconversion-based nanothermometry

Keywords: lanthanide-doped nanoparticles, upconversion, quantum yield, heterostructure, lattice strain, temperature sensing

Lanthanide (Ln 3+ )-doped upconversion nanoparticles (UCNPs) often suffer from weak luminescence intensity especially when their sizes are ultrasmall (less than 10 nm).

Enhancing the upconversion luminescence (UCL) efficiency of ultrasmall UCNPs has remained a challenge that must be undertaken if any practical applications are to be envisaged. Herein, we present a unique Ln 3+ -doped oxysulfide@fluoride core/shell heterostructure which shows efficient UCL properties under 980 nm excitation and good stability in solution. Through an epitaxial heterogeneous growth, a ~4 nm optically inert β-NaYF 4 shell coats onto ~5 nm ultrasmall Gd 2 O 2 S:20%Yb,1%Tm. These Gd 2 O 2 S:20%Yb,1%Tm@NaYF 4 core/shell UCNPs show a more than 800-fold increase in UCL intensity compared to the unprotected core, a 200fold increase in luminescence lifetime of the 3 H 4 → 3 H 6 Tm 3+ transition from 5 μs to 1000 μs and an upconversion quantum yield (UCQY) of 0.76% at an excitation power density of 155 W/cm 2 .

Likewise, Gd 2 O 2 S:20%Yb,2%Er@NaYF 4 core/shell UCNPs show a nearly 5000-fold increase of their UCL intensity compared to the Gd 2 O 2 S:20%Yb,2%Er core and a maximum UCQY of 0.61%. Further, we have shown that the I 690 /I 803 emission of Gd 2 O 2 S:20%Yb,1%Tm@NaYF 4 core/shell UCNPs is very sensitive to temperature. It is likely that this best-in-class relative thermal sensitivity of 3.9% K -1 originates from a change in the lattice strain, a property that emanates from the heterogeneous core-shell architecture. Such a result provides a proof-ofconcept for the use of these UCNPs for nanothermometry applications.

INTRODUCTION

Lanthanide (Ln 3+ )-doped upconversion nanoparticles (UCNPs) that convert low-energy infrared photons into high-energy ultraviolet (UV) or visible ones have attracted considerable attention in recent decades owing to their unique optical properties such as high photostability, long luminescence lifetimes, and sharp emission bands. [START_REF] Haase | Upconverting Nanoparticles[END_REF] These features make UCNPs extremely attractive for use as reporters in fluorescent microscopy, nanomedicine, optogenetics, and nanosensing. [START_REF] Zheng | Recent advances in upconversion nanocrystals: Expanding the kaleidoscopic toolbox for emerging applications[END_REF][START_REF] Wen | Future and challenges for hybrid upconversion nanosystems[END_REF][START_REF] Chen | Emerging Frontiers of Upconversion Nanoparticles[END_REF][START_REF] All | Expanding the Toolbox of Upconversion Nanoparticles for In Vivo Optogenetics and Neuromodulation[END_REF][START_REF] Qiu | Ratiometric upconversion nanothermometry with dual emission at the same wavelength decoded via a time-resolved technique[END_REF][START_REF] Brites | Lanthanide-Based Thermometers: At the Cutting-Edge of Luminescence Thermometry[END_REF] To date, the major bottlenecks of Ln-based UCNPs reside in the low extinction coefficients of lanthanide ions, together with low upconversion (UC) quantum yield (QY), which, in the case of very small particles, is mostly ascribed to surface quenching effects [START_REF] Zheng | Lanthanide-doped upconversion nano-bioprobes: electronic structures, optical properties, and biodetection[END_REF][START_REF] Li | Highly efficient lanthanide upconverting nanomaterials: Progresses and challenges[END_REF] ; the combination of these two factors results in low brightness. Core/shell engineering by addition of an inert shell onto the optically active core, can drastically enhance upconversion luminescence (UCL) efficiency, and has been regarded as the most efficient strategy. [START_REF] Su | The Effect of Surface Coating on Energy Migration-Mediated Upconversion[END_REF] To this end, fluoride@fluoride homogeneous (e.g. NaYF 4 :Yb 3+ ,Er 3+ @NaYF 4 11 ) or similar (e.g. NaYF 4 @NaGdF 4 [START_REF] Abel | Hard Proof of the NaYF 4 /NaGdF 4 Nanocrystal Core/Shell Structure[END_REF] ) core/shell structures have been extensively studied because fluoride host matrices are normally recognized as the most efficient family of host materials with low phonon energies for achieving highly efficient UCL. Construction of heterogeneous core/shell structures by combining two different materials can generally lead to a combination of advantages. For instance, CsPbBr 3 /NaYF 4 :Yb,Tm core/shell NPs with extended excitation wavelength to nearinfrared (NIR) and improved stability in water due to addition of the fluoride shell, have been synthesized. [START_REF] Ruan | NIR-excitable heterostructured upconversion perovskite nanodots with improved stability[END_REF] However, it still remains challenging to synthesize heterostructured nanocrystals consisting of two different families with structural and lattice mismatch.

Rare-earth oxysulfide (RE 2 O 2 S) materials such as Gd 2 O 2 S (phonon energy, ~440 cm -1 ), a family of ideal host materials for Ln 3+ dopants, have evoked much interest in various applications due to their wide band-gap, strong antioxidation, high X-ray absorption efficiency, and excellent optical performance. [START_REF] Zhang | Lanthanide-doped nanoparticles in photovoltaicsmore than just upconversion[END_REF][START_REF] Ludwig | Evaluation of Gd 2 O 2 S:Tb as a Phosphor for the Input Screen of X-Ray Image Intensifier[END_REF][START_REF] Qian | Solvothermal synthesis of columnar Gd 2 O 2 S:Eu 3+ and a comparative study with columnar Gd 2 O 3 :Eu 3+[END_REF] Specifically, Gd 2 O 2 S:Tb based phosphor has been widely used as a highly efficient X-ray scintillator material for intensifier screen, because the emission [START_REF] All | Expanding the Toolbox of Upconversion Nanoparticles for In Vivo Optogenetics and Neuromodulation[END_REF] D 4 → 7 F 5 of Tb 3+ under X-ray excitation is located in the green region which matches well with the working range of silicon-based photo-sensors. [START_REF] Machado | Xray scintillator Gd 2 O 2 S:Tb 3+ materials obtained by a rapid and cost-effective microwave-assisted solid-state synthesis[END_REF] In addition, Gd 2 O 2 S:10%Er microcrystals have shown a record-high UCQY value of 12% at 1510 nm infrared (IR) excitation, which makes them suitable for photovoltaic applications. [START_REF]Mart n-odr ue[END_REF] However, because of the limited synthesis strategies, there are no reports of for the fabrication of Gd [START_REF] Costi | Colloidal Hybrid Nanostructures: A New Type of Functional Materials[END_REF][START_REF] Carbonea | Colloidal heterostructured nanocrystals: Synthesis and growth mechanisms[END_REF] as shown in Figure 1a. A large batch (4 mmol) of Gd 2 O 2 S:20%Yb,1%Tm core-only UCNPs was first synthesized through a modified thermal decomposition protocol. The preparation of the RE 2 O 2 S@NaREF 4 core/shell NPs was then carried out via high-temperature co-precipitation as previously described [START_REF] Qian | Synthesis of Hexagonal-Phase Core-Shell NaYF 4 Nanocrystals with Tunable Upconversion Fluorescence[END_REF] .

Confirmation of the heterogeneous core/shell structure

The TEM image in Figure S1a shows good dispersity of the as-prepared Gd 2 O 2 S:20%Yb,1%Tm core-only UCNPs. The corresponding mean size of the NPs was determined to be 5.2 ± 0.9 nm (Figure S1c). The high-resolution (HR) TEM image shows several sets of clear lattice fringes in Figure S1b, indicating the fine crystallinity of the NPs. The d-spacing was measured to be 0.308 nm which can be indexed to (101) lattice plane of Gd 2 O 2 S. The TEM image in Figure 1b shows highly uniform size and morphology and the HRTEM image in Figure 1c shows good crystallinity of the Gd 2 O 2 S:20%Yb,1%Tm@NaYF 4 core/shell UCNPs. The mean size was determined to be 13.7 ± 0.8 nm in diameter (Figure S2). XRD measurements in Figure S3 confirm the pure hexagonal phase for core-only Gd 2 O 2 S and the hexagonal NaYF 4 for core/shell NPs. It should be noted that the peaks of core/shell NPs shift to the lower 2θ angle, which hints at the lattice distortion in the NPs, however, no obvious Gd 2 O 2 S diffraction patterns were observed in the XRD of core/shell. We suggest the vanishing of Gd 2 O 2 S diffraction patterns is because the very weak diffraction signals of the ultrasmall crystallites are covered by the strong diffraction of the bigger shell. A similar phenomenon has been reported in other heterogeneous α-NaREF 4 @β-NaREF 4 core/shell NPs. [START_REF] Chen | Synthesis and Luminescence Properties of Water Soluble α-NaGdF 4 /β-NaYF 4 :Yb,Er Core-Shell Nanoparticles[END_REF][START_REF] Shao | Core-mediated synthesis, growth mechanism and near-infrared lu inescence en ance ent of α-NaGdF 4 @β-NaLuF 4 :Nd 3+ core-shell nanocrystals[END_REF] To confirm the core/shell structure, a single nanoparticle was extracted from the HRTEM image in Figure 1c, as shown in Figure 1d where

Analysis of the growth kinetics

To gain insight into the growth kinetics of Gd 2 O 2 S:20%Yb,1%Tm@NaYF 4 core/shell heterostructure, we performed size and phase evolution investigation based on TEM, XRD, and SAXS characterizations. Samples were collected from the reaction solution after different heating times. TEM images in Figures 2a andS5 show the evolution of size and morphology during the reaction process. From 60 °C to 220 °C (Stage I), we observed one population of well crystallized and uniform NPs with size slightly decreasing with increasing temperature. These NPs were supposedly core-only NPs, maybe some poorly crystallized and small sized α-NaYF 4 had been removed during the washing process since it is claimed that α-NaYF 4 can be obtained below 220 °C24 . As the temperature increased, a second population of smaller NPs with a size of 2.8 nm were first observed at 250 °C. These small NPs grew slowly with rising temperature and manifestly disappeared before the 300 °C, 30 min point (end of Stage II). Meanwhile, the size of large NPs increased a little. We assume that the small NPs may be α-NaYF 4 as reported previously [START_REF] Peng | Epitaxial Growth of Highly Luminescent CdSe/CdS Core/Shell Nanocrystals with Photostability and Electronic Accessibility[END_REF][START_REF] Peng | Kinetics of II-VI and III-V Colloidal Se iconductor Nanocrystal Growt : " ocusin " of Size Distributions[END_REF][START_REF] Johnson | Self-Focusing by Ostwald Ripening: A Strategy for Layer-by-Layer Epitaxial Growth on Upconverting Nanocrystals[END_REF] , and the large NPs may be the core-only NPs with size enlargement by surface attachments of ultrasmall α-NaYF 4 . Finally, the large NPs continued to grow and small NPs disappeared as the reaction progressed (Stage III). XRD measurements further confirmed the phase evolution of the NPs. As seen in Figure S6, only the Gd 2 O 2 S phase was observed before Based on the above observations and analyses, we concluded a tentative formation process for the Gd 2 O 2 S:20%Yb,1%Tm@NaYF 4 core/shell heterostructure, as shown in Figure 2b. Firstly, cation exchange at the solid-liquid interface can contribute to lowering the interfacial energy of the system, which may form an amorphous layer. Then, with increasing temperature (Stage II), α-NaYF 4 nucleation occurs with temperature increases (Stage II). In this stage, ultrasmall α-NaYF 4 could be attached on the core-only particles by electrostatic attraction [START_REF] Chen | Photon upconversion in core-shell nanoparticles[END_REF] , or even form a thin α-NaYF 4 layer because some NPs were observed to be more round than before. Since α- Er, Tm, Yb, and Lu) were synthesized by using the same batch of core-only UCNPs. The TEM images in Figure S7 and calculated sizes in Table S3 show the slight variations in shapes and sizes of the core/shell nanoparticles, which are likely caused by the different degrees of lattice mismatch (ranging from -1.5% to 3.52%, Table S2) between the core and the shell. This establishes the feasibility of core-multishell structures based around oxysulfide cores and paves the way for layered nanocrystals with rational ion segregation.

Enhancement of luminescent properties and stability

The 4a), respectively. We observed an 836-fold enhancement in overall emission for Tm-doped UCNPs after NaYF 4 coating. As a control, we coated the core with a Gd 2 O 2 S shell using a hot injection method. As shown in Figure S8a, we obtained irregular Gd 2 O 2 S:20%Yb,1%Tm@Gd 2 O 2 S core/shell UCNPs with an average size of 14.0 ± 1.4 nm.

However, only a 20% luminescence increase was obtained for these Gd 2 O 2 S:20%Yb,1%Tm@Gd 2 O 2 S core/shell UCNPs under 980 nm excitation (Figure S8b). This confirms the advantage of a heterogeneous coating for RE 2 O 2 S UCNPs compared to Gd 2 O 2 S:20%Yb,1%Tm@Gd 2 O 2 S homogeneous core/shell UCNPs which displays sub-optimal protections by partial coating, irregular shapes, and a broad size distribution. For Gd 2 O 2 S:20%Yb,2%Er@NaYF 4 core/shell UCNPs, a nearly 5000-fold enhancement was obtained (Figure S9a). Time-resolved luminescence of both Yb 3+ /Tm 3+ -doped core-only and core/shell UCNPs were recorded. The lifetimes obtained by a monoexponential fit of the 469 nm and 803 nm emission decays were increased from 8 μs and 5 μs for the core-only (Figure 3b) to 919 μs and 997 μs for the core/shell (Figure 3c), respectively, displaying approximately 100-to 200-fold longer lifetimes. The power-dependent absolute UCQYs of the Ln 3+ -doped Gd 2 O 2 S@NaYF 4 core/shell UCNPs were measured to quantify the UCL efficiency. Figure 3d shows the maximum UCQY value of 0.76% for Gd 2 O 2 S:20%Yb,1%Tm@NaYF 4 at a power density of 155 W/cm 2 . Likewise, the maximum UCQY value of 0.61% was acquired for Gd 2 O 2 S:20%Yb,2%Er@NaYF 4 UCNPs (Figure S9b). These UCQYs values are moderate but remarkable considering the very small size of the emitting cores of these UCNPs. [START_REF] Fischer | Quantum Yields, Surface Quenching, and Passivation Efficiency for Ultrasmall Core/Shell Upconverting Nanoparticles[END_REF] In the future, preparing larger core-only NPs or screening proper shell matrices could be attempted to improve the UCQYs. Although it is claimed that RE 2 O 2 S materials have good stability [START_REF] Larquet | Metal Oxysulfides: From Bulk Compounds to Nanomaterials[END_REF] , however, there is a lack of investigation on the colloidal stability of small-sized and ligand-coated RE 2 O 2 S NPs. Stability of ligand-coated NPs in various environments (such as cyclohexane, water, and acid solution) plays a vital role in performance characterizations, construction of nanostructures and applications.

Hence, we studied the stability of ligand-coated Gd 2 O 2 S:20%Yb,1%Tm core-only and Gd 2 O 2 S:20%Yb,1%Tm@NaYF 4 core/shell UCNPs in cyclohexane and acidic ethanol solutions.

The results in Figures S10 andS11 showed that the core-only UCNPs could not be stably dispersed in cyclohexane at room temperature for a long time without aggregation and they showed low resistance to acidic conditions (acidified ethanol). Comparatively, the core/shell

UCNPs with an NaYF 4 shell showed good stability in both cyclohexane and acid solution, proving the enhanced stability after coating a fluoride shell.

Temperature effect on the luminescence spectra

Lattice strain in the heterostructure can tune the optical properties of the nanocrystals and can be sensitive to various stimuli such as external mechanical forces and temperature. To verify this, we first implemented temperature-dependent UCL spectra measurements for both Yb 3+ /Tm 3+ and Yb 3+ /Er 3+ -doped Gd 2 O 2 S@NaYF 4 UCNPs. Figure 4a and Figure S12 show that the UCL intensity increases continuously with increasing temperature from 293 K to 473 K. The intensities of the Tm 3+ bands at 469 nm, 640 nm, 690 nm ( 3 F 2,3 → 3 H 6 ), and 803 nm were 3, 2, 10, and 3 times higher at 473 K than those at 293 K, respectively (Figure 4b). Notably, the reason for the greater enhancement of the 690 nm band is that 3 H 4 and 3 F 2,3 energy levels are thermallycoupled, thus favoring the population at the higher level of 3 F 2,3 at an elevated temperature. This indicates that the UCNPs have a high sensitivity to temperature and thus can be good candidates as thermometers. A similar thermal enhancement was observed for the UCL of Yb 3+ In recent years, thermally enhanced UCL has attracted a lot of attention and several hypotheses such as surface phonon assistance [START_REF] Zhou | Activation of the surface dark-layer to enhance upconversion in a thermal field[END_REF] , surface molecule desorption [START_REF] Hu | Energy Loss Mechanism of Upconversion Core/Shell Nanocrystals[END_REF] , and energy transfer suppression induced by lattice expansion [START_REF] Cui | Size-dependent abnormal thermo-enhanced luminescence of ytterbium-doped nanoparticles[END_REF] have been proposed. In our case, the shell thickness of small (< 30 nm) Gd 2 O 2 S@NaYF 4 core/shell UCNPs is only about 4 nm which cannot completely prevent surface quenching by water molecules. Hence, we thought that the UCL enhancement with increasing temperature in our UCNPs was caused by water desorption [START_REF] Chen | Emerging Frontiers of Upconversion Nanoparticles[END_REF][START_REF] Hu | Energy Loss Mechanism of Upconversion Core/Shell Nanocrystals[END_REF][START_REF] Shi | Thermal enhancement of upconversion emission in nanocrystals: a comprehensive summary[END_REF] . To confirm this hypothesis, we carried out heating-cooling cycle experiments for Gd 2 O 2 S:20%Yb,2%Er@NaYF 4 core/shell UCNPs. As shown in Figure S14, in the first 2 cycles, thermal UCL enhancement was observed as the temperature increased. However, the UCL intensity continued to increase in the cooling process, which could be due to water desorption. In the last 3 cycles, the UCL intensity at 473 K was always higher than that at 293 K, which means surface attached water molecules were almost removed and thermal quenching was dominant.

temperature. Figure 4e shows that the ratio first increases linearly from 293 K to 423 K, then remains almost stable from 423 K to 463 K, and finally drops at 473 K. The ratio at 423 K is heating and differential lattice expansion. An increasing temperature can enhance the non- the lattice distortion will vary from tensile strain to compressive strain due to large differences in thermal expansion rates of plane spacings between core and shell materials.

radiative relaxation of the 4 S 3/2 → 4 F 9/2 transition, leading to an increase of the I 652 /I 540 ratio. This can, however, only explain the change in the first temperature range from 293 K to 423 K. As the temperature continues to rise, the I 652 /I 540 ratio nearly remains constant and even drops at 473 K.

Hence, lattice expansion seems to be involved as well. We first carried out in-situ XRD measurements at different temperatures (from 323 K to 573 K) for both core-only and core/shell UCNPs, as shown in Figures S15 andS16. To ensure the reversibility of the lattice constants after heating process, the XRD patterns of the UCNPs were measured twice at 323 K. Note the Gd 2 O 2 S:Yb,Tm core-only UCNPs were synthesized by a solid state method to avoid surface ligands and used as a reference [START_REF] Osseni | Gadolinium oxysulfide nanoparticles as multimodal imaging agents for T 2 -weighted MR, X-ray tomography and photoluminescence[END_REF] , and the surface ligand of Gd 2 O 2 S:Yb,Tm@NaYF 4 core/shell UCNPs were removed by an acid treatment method (see Supporting Information). The lattice constants were determined by Pawley method fitting [START_REF] Pawley | Unit-cell refinement from powder diffraction scans[END_REF] , then the cell volumes at different temperatures were calculated (see Tables S4 andS5). Further, the linear fitting of cell volume versus temperature plots revealed thermal volume expansion rates of 0.007 Å 3 /K and 0.021 Å 3 /K for Gd 2 O 2 S:Yb,Tm core-only and Gd 2 O 2 S:Yb,Tm@NaYF 4 core/shell, respectively (Figure 4f).

Based on the previous XRD and HRTEM analyses, existence of tensile strain in the NaYF 4 shell has been confirmed at ambient temperature. Therefore, we could postulate the variation of the interface lattice distortion with elevating temperature as illustrated in Figure 4g. The lattice distortion decreases with elevating temperature because of the large difference of the expansion rate between core and shell (shell expansion rate is 300% higher than core) and reach an equilibrium. This process has recently been reported as lattice self-adaptation [START_REF] Wu | Nanoscale Ultrasensitive Temperature Sensing Based on Upconversion Nanoparticles with Lattice Self-Adaptation[END_REF] . It is inferred that further increasing the temperature would generate compressive strain in the shell. The varied strain in the shell generates opposite strain in the core, producing variation in the local structure in the interfacial area, further changing the optical properties of such heterostructures. [START_REF] Bao | Interfacial Stress-Modulated Mechanosensitive Upconversion Luminescence of NaErF 4 Based Heteroepitaxial Core-Shell Nanoparticles[END_REF] When the temperature increases from 293 to 423 K, decrease of compressive strain in the core leads to enhancement of 4 S 3/2 → 4 F 9/2 non-radiative transition, linearly increasing the luminescence intensity ratio I 652 /I 540 , which is in a good agreement with investigations of the Er 3+ emission of other NaYF 4 based heterostructures subjected to heating [START_REF] Zhao | Accurate Control of Core-Shell Upconversion Nanoparticles through Anisotropic Strain Engineering[END_REF] and mechanical forces [START_REF] Lay | Mechanosensitive Upconversion with Cubic-P ase Heteroepitaxial Core-S ell Nanoparticles[END_REF][START_REF] Lay | Upconverting Nanoparticles as Optical Sensors of Nano-to MicroNewton Forces[END_REF] . Then, the I 652 /I 540 ratio remains nearly constant in the range of 423~463 K due to zero strain at the interfacial area. Finally, the inverse tensile strain in the core decreases the I 652 /I 540 ratio at 473 K.

Importantly, when the lattice strain at the interface has a strong influence on the population of thermally-coupled emissive states, it can be harnessed to enhance the sensitivity to temperature. [START_REF] Wu | Nanoscale Ultrasensitive Temperature Sensing Based on Upconversion Nanoparticles with Lattice Self-Adaptation[END_REF][START_REF] Zhao | Accurate Control of Core-Shell Upconversion Nanoparticles through Anisotropic Strain Engineering[END_REF][START_REF] Lay | Mechanosensitive Upconversion with Cubic-P ase Heteroepitaxial Core-S ell Nanoparticles[END_REF]

Towards a nanothermometry probe

As a proof-of-concept, we demonstrated a high response of UCL to temperature change using our custom-designed Ln 3+ doped Gd 2 O 2 S@NaREF 4 heterostructures. Since 980 nm excited 3 F 2,3 UCL (arising from their respective thermally coupled energy levels) comply with Boltzmann distribution theory [START_REF] Wade | Fluorescence intensity ratio technique for optical fiber point temperature sensing[END_REF][START_REF] Bradac | Optical Nanoscale Thermometry: From Fundamental Mechanisms to Emerging Practical Applications[END_REF] , the luminescence intensity ratio (LIR) can be expressed following the equation:

(

where I 2 and I 1 corresponds to the integrated fluorescence intensities of high-energy level and low-energy level, respectively; C is a constant; △E is the energy gap; k is the Boltzmann constant and T is the absolute temperature. The obtained LIR and Ln(LIR) values were fitted with Eq. ( 1)

for Yb 3+ /Tm 3+ and Yb 3+ /Er 3+ doped Gd 2 O 2 S@NaYF 4 core/shell UCNPs respectively, as shown in Figures 5a and5d (the fitting results can be found in the Supporting Information).

To further evaluate the temperature sensing behavior of our UCNPs, the absolute thermal sensitivity (S a ) and the relative temperature sensitivity (S r ) were calculated using the following equations [START_REF] Wade | Fluorescence intensity ratio technique for optical fiber point temperature sensing[END_REF][START_REF] Bradac | Optical Nanoscale Thermometry: From Fundamental Mechanisms to Emerging Practical Applications[END_REF] :

(

(3)

The S a and S r values for both Yb 3+ /Tm 3+ and Yb 3+ /Er 3+ doped UCNPs have been presented in Figures 5b and5e. As can be seen, the Yb 3+ /Tm 3+ doped UCNPs exhibit the maximum S a at 473 K with 0.0069% K -1 and S r at 293 K with 3.9% K -1 ; the Yb 3+ /Er 3+ doped UCNPs exhibit the maximum S a at 473 K with 0.4% K -1 and S r at 293 K with 1.1% K -1 . We performed the same method to assess the temperature sensing capacity for Gd 2 O 2 S:20%Yb,1%Tm@NaLuF 4 UCNPs which has a different shell coating matrix, as shown in Figure S17. The maximum S a at 473 K with 3.0% K -1 and S r at 293 K with 3.3% K -1 were obtained for the UCNPs. To assess the photochemical stability of our luminescent nanoprobes, we performed 5 heating-cooling cycles for both Yb 3+ /Tm 3+ and Yb 3+ /Er 3+ doped UCNPs. According to the data in Figure 5c and 5f, the reproducibility value (R) was calculated to be higher than 95% by the method reported previously [START_REF] Brites | Lanthanide-Based Thermometers: At the Cutting-Edge of Luminescence Thermometry[END_REF] .

Table 1 summarizes the sensitivities obtained from other previously reported material systems using the same intensity ratios of Tm 3+ and Er 3+ for comparison. It can be found that the maximum S r of our Gd 2 O 2 S:20%Yb,1%Tm@NaYF 4 core/shell UCNPs is the highest amongst Yb 3+ /Tm 3+ doped systems, and even the NaLuF 4 -coating UCNPs have a greater S r value than previously reported Lu-based materials in the same temperature range reported in the literature [START_REF] Min | Temperature sensing based on the up-conversion emission of Tm 3+ in a single KLuF 4 microcrystal[END_REF] .

Also, the Yb 3+ /Er 3+ doped Gd 2 O 2 S@NaYF 4 core/shell UCNPs have an excellent maximum S r value which is comparable with the values found for other systems. The increased performance of the thermal sensitivity of our UCNPs seems to be associated with interfacial lattice strain that can enhance the temperature sensitivity. Note that all the compositions of our NPs in this study are still not optimized towards temperature sensitivity, which further highlights the potential of Ln 3+ doped RE 2 O 2 S@NaREF 4 as a new kind of luminescence-based nanothermometer.

CONCLUSIONS

In conclusion, we have designed an upconversion core/shell heterostructure by combining RE 2 O 2 S and NaREF 4 matrices formed by heterogeneous epitaxial deposition mechanism as demonstrated by growth kinetic studies. Upon addition of inert NaREF 4 shells, we observed remarkable improvements of the UCL intensities, lifetimes, and UCQYs, as well as stability for the core/shell UCNPs compared to those features of the cores. These results confirmed the effective protection brought by the shell to the core not only in the UCL properties but also the enhanced stability in different solvent environments. The distortion at the core/shell interface induced by lattice mismatch was observed by HRTEM and XRD analyses. For both Yb 3+ /Er 3+ and Yb 3+ /Tm 3+ doped Gd 2 O 2 S@NaYF 4 core/shell UCNPs, the temperature-induced lattice distortion provides the ability to design sensitive optical probes response to temperature and mechanical forces. In addition, tunable interfacial strain by varying shell components may give possibilities to further improve the sensitivities. The Yb 3+ /Tm 3+ doped Gd 2 O 2 S@NaYF 4

Figure 1 .

 1 Figure 1. Structure and morphology. (a) Concept graph of Gd 2 O 2 S@NaYF 4 core/shell heterostructure. (b, c) Low-and high-resolution TEM images of Gd 2 O 2 S:20%Yb,1%Tm@NaYF 4 core/shell UCNPs, respectively. (d) HRTEM image of a single nanoparticle extracted from red rectangle of (c), corresponding FFT (top right) and IFFT (bottom right) diffraction patterns showing core (red dashed line) and shell (green dashed line) regions in the nanoparticle. (e, f) ABF and HAADF-STEM images of the core/shell nanoparticles, respectively. (g) Line scan going across a single nanoparticle and the resulting EELS profiles of Gd and F. (h) Element mapping of O, F, S, Y, Gd, Tm, and Yb showing the elemental distribution.

  Figure 1h confirm the composition and the distribution of elements in the

Figure 2 .

 2 Figure 2. Investigation of growth kinetics. (a) TEM images showing status of nanoparticles at different stages, and sizes evolution during the synthesis. The sizes were statistically measured on at least 50 particles. (b) Proposed formation processes for RE 2 O 2 S@NaYF 4 core/shell heterostructure. (c) SAXS curves for nanoparticles at 60 °C and 300 °C with heating times from 0 to 60 min. (d) Rg and Porod volume extracted from the SAXS data of (c).

NaYF 4

 4 is the kinetic product, it dissolves quickly at elevated temperature and nucleates into the thermodynamic β-NaYF 4 phase which then deposits on the larger RE 2 O 2 S NPs (Stage III). Note that this dissolution/deposition process is very rapid. Driven by Ostwald ripening, the larger NPs continually increased in size.Additionally, SAXS measurements were carried out to investigate global dimensions of the NPs in different reaction phases. As shown in Figure2c, the superimposition of the SAXS curves shows an increase of the SAXS signal during the temperature rise phase (60 °C vs. 300°C, 0 min), suggesting an increase of the concentration of α-NaYF 4 NPs. From 0 to 30 min at 300°C, the changes observed on the SAXS curves highlight a phase growth of nanoparticle size. Next, from 45 to 60 min at 300°C, β-NaYF 4 phase occurs and grows as demonstrated by the formation of new peaks in the SAXS curves. We noticed that the remarkable modulations in the curves indicate a good monodispersity of the particles in solution. By extracting the data from the SAXS curves, the radius of gyration (Rg) and the Porod volume of particles were determined to follow the particle growth (Figure2d). It can be seen that the global size increases a little during the temperature rising phase, which is consistent with observation by TEM. In addition, the variations of Rg and volume of particles during the synthesis are in good agreement with the observations by TEM.The successful preparation of RE 2 O 2 S@NaYF 4 core/shell heterostructure encouraged us to synthesize other heterostructures with different RE-shells. Accordingly, a series of Gd 2 O 2 S:20%Yb,1%Tm@NaREF 4 core/shell NPs (RE = La, Ce, Pr, Nd, Sm, Eu, Gd, Tb, Dy, Ho,

  Figure 3. UCL properties. (a) UCL spectra of Gd 2 O 2 S:20%Yb,1%Tm core-only and Gd 2 O 2 S:20%Yb,1%Tm@NaYF 4 core/shell UCNPs under 980 nm laser excitation. (b, c) UCL decays of core-only and core/shell UCNPs, respectively. (d) Absolutely determined UCQYs of Gd 2 O 2 S:20%Yb,1%Tm@NaYF 4 core/shell UCNPs under different power densities.

Figure 4 .

 4 Figure 4. Analyses of temperature-dependent UCL spectra. (a, b) Variation of UCL spectra and integrated intensities of the peaks of Tm 3+ , respectively, of Gd 2 O 2 S:20%Yb,1%Tm@NaYF 4 core/shell UCNPs with increasing temperature. Inset energy level diagram shows the upconversion mechanism of Tm 3+ ions. (c, d) Variation of UCL spectra and integrated intensities of the peaks of Er 3+ , respectively, of Gd 2 O 2 S:20%Yb,2%Er@NaYF 4 core/shell UCNPs with increasing temperature. Inset energy level diagram shows the upconversion mechanism of Er 3+ ions. (e) Temperature-dependent red/green emission intensity ratios (I 652 /I 540 ) of Gd 2 O 2 S:20%Yb,2%Er@NaYF 4 core/shell UCNPs. The red straight line is plotted by linear fitting in the temperature range of 293~423 K. (f) Cell volume as a function of temperature plots for both core and core/shell compositions. The thermal volume expansion rates can be determined by linearly fitting the respective data points. (g) Schematic illustration of lattice distortion variation with increasing temperature. It shows that

→ 3 H 6 (

 36 I 690 ) and 3 H 4 → 3 H 6 (I 803 ) of Tm 3+ , 2 H 11/2 → 4 I 15/2 (I 520 ) and 4 S 3/2 → 4 I 15/2 (I 540 ) of Er 3+

Figure 5 .

 5 Figure 5. Thermal sensitivity. (a) LIR vs T, (b) dependence of absolute (S a ) and relative (S r ) thermal sensitivity on temperature, and (c) heating-cooling cycles for I 690 /I 803 of Gd 2 O 2 S:20%Yb,1%Tm@NaYF 4 core/shell UCNPs. (d) Ln(LIR) vs 1/T, (e) dependence of S a and relative S r on temperature, and (f) heating-cooling cycles for I 520 /I 540 of Gd 2 O 2 S:20%Yb,2%Er@NaYF 4 core/shell UCNPs.

  Both Gd 2 O 2 S and NaYF 4 are known to have their thermodynamically stable products in a hexagonal lattice, which inspired us to construct a heterogeneous structure Gd 2 O 2 S@NaYF 4 . We first theoretically investigated the lattice mismatch between hexagonal RE 2 O 2 S and NaREF 4 (RE = Y, La, Ce, Pr, Nd, Sm, Eu, Gd, Tb, Dy, Ho, Er, Tm, Yb, and Lu). TableS1shows the identical d-spacings of (101) or (011) planes of RE 2 O 2 S and (110) or (101) planes of NaREF 4 , which were extracted from their respective XRD files. The lattice mismatch values were further calculated and given in TableS2. The Gd 2 O 2 S and NaYF 4 lattices show only a mismatch of 0.5%, indicating the feasibility of our concept of Gd 2 O 2 S@NaYF 4 heterostructure formed by epitaxial growth

	RESULTS AND DISCUSSION
	Design and synthesis

2 O 2 S based heterogeneous core/shell UCNPs with efficient UCL, nor for their application as nanothermometers.

Herein, we present a strategy to construct Ln 3+ -doped Gd 2 O 2 S@NaREF 4 heterogeneous core/shell structures. An inert NaYF 4 shell of ~4 nm thickness was successfully deposited on the ultrasmall (~5 nm) Gd 2 O 2 S:20%Yb,1%Tm core-only UCNPs. The growth kinetics of RE 2 O 2 S@NaREF 4 core/shell NPs were studied by meticulously analyzing TEM, XRD, and small-angle X-ray scattering (SAXS) in different reaction phases. UCL properties including emission intensity, lifetime, and UCQY of Ln 3+ -doped Gd 2 O 2 S@NaYF 4 core/shell UCNPs were systematically investigated. Stability of the core-only and core/shell UCNPs in different environments was monitored. Moreover, temperature-dependent UCL spectra and in-situ XRD measurements manifest a correlation between selected emission bands ratio of Er 3+ and the lattice strain. Finally, we show that Ln 3+ -doped Gd 2 O 2 S@NaREF 4 core/shell UCNPs have a great capacity for nanothermometry application with highest-in-class thermal sensitivities.

  inert NaYF 4 shell protects the Ln 3+ -doped Gd 2 O 2 S cores from surface quenching, thus significantly boosting the UCL properties. As shown in Figure3a, under illumination with a continuous wavelength 980 nm laser, Gd 2 O 2 S:20%Yb,1%Tm core-only and Gd 2 O 2 S:20%Yb,1%Tm@NaYF 4 core/shell UCNPs exhibit emission bands at 469 nm, 640 nm, and 803 nm ascribed to the 1 G 4 → 3 H 6 , 1 G 4 → 3 F 4 , and 3 H 4 → 3 H 6 transitions of Tm 3+ (see energy diagram in Figure

  /Er 3+ -doped Gd 2 O 2 S@NaYF 4 UCNPs, shown in Figures4c, S13, and 4d. As the temperature increases, the intensities of the Er 3+ peaks at 520 nm ( 2 H 11/2 → 4 I 15/2 ), 540 nm ( 4 S 3/2 → 4 I 15/2 ), and 652 nm ( 4 F 9/2 → 4 I 15/2 ) increase by factors of 3.9, 1.1, and 2.8 at 473 K compared to those at 293 K.

Table 1 .

 1 Comparison of the relative thermal sensitivities (S r ) of Yb/Tm or Yb/Er doped systems in the similar temperature range. Luminescence intensity ratios (LIR) from the transitions 3 F 2,3 → 3 H 6 / 3 H 4 → 3 H 6 for Tm 3+ and 2 H 11/2 → 4 I 15/2 / 4 S 3/2 → 4 I 15/2 for Er 3+ under 980 nm laser excitation. Values determined by us for this study, according to the data published by the authors in the original literature.

	Material	T-range (K)	FIR	Sr (% K -1 )	Ref.
	Gd 2 O 2 S:Yb,Tm@NaYF 4	293-473	I 690 /I 803	3.9 at 293 K	This work
	Gd 2 O 2 S:Yb,Tm@NaLuF 4	293-473	I 690 /I 803	3.3 at 293 K	This work
	GdYO 4 :Yb,Tm@SiO 2	298-333	I 700 /I 800	3.3 at 298 K	46

°C but the peaks broadened with increasing temperature due to the formation of smaller size of NPs and a very thin amorphous layer on the surface of the core via the cation exchange process between the core and the precursor solution. This phenomenon is often observed in the synthesis of heterogeneous nanostructures.[START_REF] Zhao | Growth of hexagonal phase sodium rare earth tetrafluorides induced by heterogeneous cubic phase core[END_REF] Subsequently, intense XRD signals of α-NaYF 4 phase appeared at 250 °C and overshadowed the intensity of Gd 2 O 2 S, indicating plenty of α-NaYF 4 produced by burst nucleation. With the disappearance of α-NaYF 4 , the peaks of the β-NaYF 4 phase became dominant (FigureS6b).

165% higher than that at 293 K. As we know, the variation of 4 S 3/2 → 4 F 9/2 non-radiative

ACKNOWLEDGMENTS

We are grateful to Dr. W. Zheng, Dr. M. Sliwa, Dr. C. Würth, and Dr J-C Micheau for the fruitful discussions. Q. Zou acknowledges the support from the French Research ministry in the form of a PhD scholarship and the DAAD short-term grant from German government. X. Chen acknowledges the support from the NSFC (No. U1805252) and the CAS/SAFEA International Partnership Program for Creative Research Teams. We thank the federation FERMAT of Toulouse for giving us access to the SAXS bench.

Apart from the UCL enhancement, the intensity ratio of I 652 /I 540 changed as a function of relaxation can influence the I 652 /I 540 ratios. Here two factors should be taken into account: nanoprobes exhibits a relative thermal sensitivity of 3.9% K -1 , the highest reported yet in this range of temperature, to the best of our knowledge. Taken together, our findings provide a new

Notes

The authors declare no competing financial interest.


Alberto José Fernández-Carrión

Julie Bourret

James Sharp

Mengchuan Xing

Jean-René Duclère

Maggy Colas

Julie Cornette

Pierre Carles

Cécile Genevois

François Brisset

Sébastien Chenu

Mathieu Allix

Gaëlle Delaizir
email: gaelle.delaizir@unilim.fr

Elaboration and luminescence properties of Laser Heated Pedestal Growth (LHPG) Sr3Al2O6-based fibers

INTRODUCTION

Despite the successful applications of silica-based glass optical fibers as sensors or fiber lasers, they are not competitive for applications requiring extreme environmental conditions such as high temperatures, radioactivity, corrosion or high energy fluences. Laser-heated pedestal growth (LHPG) and micro-pull-down techniques have been well studied and used with increasing success to prepare crystalline fibers that are thermally, mechanically and chemically more robust than glass fibers. [START_REF] Wang | Few-Mode Cr-Doped Crystalline Core Fibers for Fiber Amplifier[END_REF][START_REF] Feigelson | Pulling optical fibers[END_REF][START_REF] Yoon | Crystal growth of dislocation-free LiNbO3 single crystals by micro pulling down method[END_REF][START_REF] Yoshikawa | Challenge and study for developing of novel single crystalline optical materials using micro-pulling-down method[END_REF][START_REF] Andreeta | In Springer Handbook of Crystal Growth[END_REF] The former is basically a miniature floating-zone-like method where the heating element is replaced by a focused laser ring to generate the molten zone. This can achieve rather high temperatures, and thereby highly refractive compositions can be melted in. The LHPG technique also has the advantage of being a containerless method and thereby preventing contamination from the crucible, which is of utmost importance as high quality optical materials are pursued.

The LHPG technique has been employed to prepare several single crystal or polycrystalline optical fibers based on binary, ternary or even quaternary compounds. [START_REF] Nie | Growth of single-crystal YAG fiber optics[END_REF][START_REF] Shi | High-concentration Er:YAG single-crystal fibers grown by laser-heated pedestal growth technique[END_REF][START_REF] Nubling | Single-crystal laser-heated pedestal-growth sapphire fibers for Er:YAG laser power delivery[END_REF][START_REF] Jundt | Characterization of single-crystal sapphire fibers for optical power delivery systems[END_REF][START_REF] Rey-García | Laser floating zone growth of Yb, or Nd, doped (Lu0.3Gd0.7)2SiO5 oxyorthosilicate single-crystal rods with efficient laser performance[END_REF][START_REF] Assi | Towards good quality Bi2ZnB2O7 fibers grown by the micro-pulling down technique[END_REF][START_REF] Alimov | Growth and Characterization of Neodymium-Doped Yttrium Scandate Crystal Fiber with a Bixbyite-type Crystal Structure[END_REF] Among them, aluminate based-compositions such as YAG or sapphire have proven high optical efficiency and are attractive for optical applications. [START_REF] Nie | Growth of single-crystal YAG fiber optics[END_REF][START_REF] Shi | High-concentration Er:YAG single-crystal fibers grown by laser-heated pedestal growth technique[END_REF][START_REF] Nubling | Single-crystal laser-heated pedestal-growth sapphire fibers for Er:YAG laser power delivery[END_REF][START_REF] Jundt | Characterization of single-crystal sapphire fibers for optical power delivery systems[END_REF][START_REF] Bera | Dopant segregation in YAG single crystal fibers grown by the laser heated pedestal growth technique[END_REF] Nevertheless, even when the usefulness of the LHPG technique has been proved, it has been scarcely explored so far. The main compositional requirement is the existence of a congruent crystallization from the molten state.

Highly transparent polycrystalline ceramic, based on Sr3Al2O6 composition has been fabricated by full crystallization from the parent glass. [START_REF] Alahraché | Perfectly Transparent Sr3Al2O6 Polycrystalline Ceramic Elaborated from Glass Crystallization[END_REF] This compound was crystallized from a bulk-glass precursor prepared by the innovative aerodynamic levitation method and features optical transparency both in the visible and infrared (until 6 m) spectral ranges. The authors attributed its high transparency to the high density (i.e. complete absence of porosity) and micrometer-scale grains with very thin grain boundaries, as well as its intrinsic cubic crystalline structure (S.G. Pa-3) leading to optical isotropy (i.e. no birefringence is present). This composition is well adapted for laser applications as illustrated by the luminescence spectra of Sr3Al2O6 ceramics doped with various rare-earth ions. [START_REF] Alahraché | Perfectly Transparent Sr3Al2O6 Polycrystalline Ceramic Elaborated from Glass Crystallization[END_REF][START_REF] Singh | Radiative properties of Er 3+ doped and Er 3+ /Yb 3+ co-doped Sr3Al2O6 phosphors: exploring the usefulness as a phosphor material[END_REF] The main drawback of the aerodynamic levitation method is the limited size of the samples, typically small beads of diameter 3-5 mm.

Owing to the crystal structure features, chemical stability, and optical properties, the fabrication of an active Sr3Al2O6 transparent ceramic fiber may be a significant breakthrough for laser-based technologies. Nevertheless, the optical transparency of this composition highly depends on the container employed during preparation. Indeed, Al Saghir et al. demonstrated that transparency cannot be attained by using the classic melt-quenching process using Pt/Rh crucibles owing to dark opaque materials being obtained. [START_REF] Al Saghir | Céramiques transparentes par cristallisation complète du verre : application aux aluminosilicates de strontium[END_REF] This drawback may be overcome by using the LHPG method as it eliminates contamination problems arising from crucibles and furnace components. [START_REF] Feigelson | Pulling optical fibers[END_REF][START_REF] Andreeta | In Springer Handbook of Crystal Growth[END_REF] In this work, we present a novel process to prepare the green-bodies (feed pedestals) by using a microextrusion machine to extrude the source rods for the growth of Sr3Al2O6 and Er 3+doped Sr3Al2O6 crystalline fibers by the LHPG technique. In addition, we employed two different kinds of seeds: Sr3Al2O6 glass and transparent ceramic seeds in place of the more commonly used crystallographic seed. A careful look at the structure, microstructure and the photoluminescence (PL) of Er-doped crystalline fibers is presented and compared to Sr3Al2O6 glass and polycrystalline bulks. The high crystal perfection obtained results in very intense and better defined emission bands when compared with the polycrystalline material counterpart.

MATERIALS AND METHODS

Preparation of Sr3Al2O6 source rods for LHPG experiment

Source rods were prepared by mixing Al2O3 (99.7%, Sigma-Aldrich) with 1 wt.% of dispersant in an agate mortar. Then, it was added the appropriate molar ratio of SrCO3 (99.9%, sigma Aldrich) and Er2O3 along with 40 % of the total weight of methyl cellulose as binder and 3 vol.% of glycerol as lubricant. The previous mixture was mixed in an agate mortar until producing a homogeneous paste which was introduced three times into a three rolls mill system to eliminate agglomerates. The resulting paste was extruded using homemade extrusion equipment through a Ø 1 mm stainless-steel die (Figure S1). The extruded green rods were air dried overnight and then used for drawing the fibers. The manual extrusion process led to shorter pedestals, with irregular shapes and superficial defects.

Preparation of Sr3Al2O6 glass and transparent ceramic used as seed

The Er-doped Sr3Al2O6 glasses were synthesized from high purity SrCO3 (99.99%), Al2O3 (99.98%) and Er2O3 (1 at% Er 3+ substitutes 1 at% Sr 2+ ) oxides using an aerodynamic levitator equipped with two CO2 lasers. Er2O3 was calcined at 900 °C overnight. The starting materials were first weighed in stoichiometric amounts, mixed, and pressed into pellets. To elaborate a glass bead, a piece of the pellet was placed in the levitator system's conical nozzle and melted, at around 1800 °C, using the CO2 lasers. The melt was levitated by an argon flow, and glass beads (1.5 mm in diameter) were obtained by quenching (realized by cutting off the two laser beams with a cooling rate of approximately 300°C/s). The corresponding transparent ceramic was obtained by congruent and full crystallization from the parent glass by a heat treatment at 840 °C for 5 h in an open-air tubular furnace. [START_REF] Alahraché | Perfectly Transparent Sr3Al2O6 Polycrystalline Ceramic Elaborated from Glass Crystallization[END_REF] These Sr3Al2O6 glasses and transparent ceramics were used as seeds by cutting the bulk materials into parallelepipeds with size 1 x 1 x 5 mm and used as a reference to compare the photoluminescence (PL) properties of the Sr3Al2O6 fiber grown by the LHPG technique.

Fiber preparation

The fibers were obtained by an LHPG equipment. Micro-extruded rods were used as source materials and Sr3Al2O6 glass or polycrystalline ceramic were used as seeds. The fibers were grown at 0.14 mm/min and common incident CO2 laser power of approximately 8 W was maintained during growth. Symmetrical heating of the melt was obtained using a standard reflaxicon arrangement. Fibers were grown 'open loop', i.e. without active diameter control. The low density of the source rods made fiber growth difficult to stabilise and hence the relatively slow pull rate of the fibers. The final fibers were approximately 800 µm in diameter with lengths ranging from 5 to 12 mm.

Characterizations

Electron backscatter diffraction (EBSD) maps were recorded on the Sr3Al2O6 LHPG fibers using an OIM TSL/EDAX system mounted on a FEG-SEM (Zeiss SUPRA 55 VP) system in order to confirm the single or polycrystalline nature of our fibers. Laboratory X-Ray Powder Diffraction (XRD) measurements were performed on crushed powder fibers using a Bragg-Brentano D8 Advance Bruker laboratory diffractometer (Cu Kα radiation) equipped with a lynxEye XE detector.

The nanostructure of the fiber was characterized by Transmission Electron Microscopy (TEM) micrographs, scanning transmission electron microscopyhigh angle annular dark field (STEM-HAADF) imaging and EDS elemental analyses. These experiments were performed on a JEOL ARM 200F (JEOL Ltd.) Cold FEG TEM operating at 200kV, equipped with a double spherical aberration corrector and fitted with a JEOL SDD CENTURIO EDS system. STEM-HAADF imaging and STEM-EDS elemental maps were acquired with a 0.1nm probe size. The samples were prepared prior to (S)TEM observations by Focused Ion Beam (FIB, Zeiss Crossbeam 550).

Raman spectroscopy characterizations of the LHPG fiber were performed using an In Via Reflex Renishaw Raman spectrophotometer. The spectra were recorded using a x50 objective with 532 nm wavelength excitation and low power at the sample (3 mW) to avoid any damage. Raman maps have been recorded with high-resolution mode, using a 2400 grooves/mm grating. An area of 58 x 66 m 2 has been studied and maps have been reconstructed using principal component analysis (PCA) method on a collection of 1020 spectra. Data have been truncated between 668 and 810 cm -1 and the baseline subtracted. The first two components explain 96.7% of the data.

The microstructure of the LHPG fibers and the chemical compositions were observed and checked using scanning electron microscopy (EDS-SEM, JEOL IT 300 LV).

Room temperature Near InfraRed photoluminescence (PL) properties were measured using a Horiba-Jobin-Yvon Fluorolog 3 spectrofluorimeter, operated in reflective geometry.

For the recorded emission spectra in the 1400-1650 nm range, the data step was fixed to 0.5 nm, with a 2.5 nm emission slit size which was found appropriate for comparing all the analyzed samples. For measuring the PL decay curves, an initial delay time of 0.2 ms was set. Data were then acquired with a time interval step of 0.2 ms, up to a final decay time of 50 ms. Lifetime values were extracted through a fitting procedure of the PL decay curves data, which employs either some single or double exponential decay function(s). Average lifetime values are finally calculated.

RESULTS AND DISCUSSION

Fiber structure

Sr3Al2O6 fibers were successfully obtained from extruded source rods of glass or transparent ceramic used as seeds as shown in Figure 1, although the drawing process was easier when extruded transparent seed ceramic were employed. It should be stressed that the drawing of Sr3Al2O6 fibers is highly dependent on the initial state of the pedestals. Initially, sintered square pedestals of 1.5mm cross section were employed. Even when small fibers could be grown, their lengths were typically shorter than 9 mm (Figure S1) and the cross section of the fibers were difficult to keep constant. The effect of the pedestal shape was discussed by Reyes Ardila et al. [START_REF] Reyes Ardila | Single-crystal SrTiO3 fiber grown by laser heated pedestal growth method: influence of ceramic feed rod preparation in fiber quality[END_REF] for the preparation of SrTiO3 fibers by LHPG. They found that rod pedestals prepared by extrusion were more suitable than squared-shape pedestals because the laser beam matches up better with the feed rod geometry. The authors also indicated some drawbacks arising from the extrusion process. Among them, lamination, surface and edge tearing. For this reason, we incorporated several steps when preparing the green bodies source rods. These include the addition of glycerol as lubricant, as well as agglomerate elimination through a three rolls mill system. Then, the green bodies were extruded by using 3D-ceramic technology which allows a constant pressure during the extrusion process, leading to perfectly straight pedestal and enhancing the final quality of the source rods. The as-prepared rods were dried in air, and further annealing treatments were avoided in order to limit furnace contamination, as well as the shrinkage of the pedestal which led to pucker them. ). Rietveld refinement was employed to estimate the percentage of this impurity which is about 2.7 vol%. We believe that the formation of this phase is related to some Al2O3 volatilization during the fiber drawing process. A similar phenomenon has been observed in Al2O3 micropulling down by Ghezal et al. [START_REF] Ghezal | Effect of Pulling Rate on Bubbles Distribution in Sapphire Crystals Grown by the Micropulling Down (μ-PD) Technique[END_REF] This hypothesis is supported by the fact that the raw rod materials can be crystallized as a Sr3Al2O6 single phase upon heating at 1100 °C for 6 hours in conventional furnace (Figure S2).

Although the average crystal structures of Sr3Al2O6 and Sr10Al6O19 present different symmetry, both can be described in terms of perovskite-based superstructures (Figure 2c-d). [START_REF] Alonso | Tristrontium dialuminum hexaoxide: an intricate superstructure of perovskite[END_REF][START_REF] Kahlenberg | Synthesis and crystal structure of Sr10Al6O19: a derivative of the perovskite structure type in the system SrO-Al2O3[END_REF] On the one hand, Sr3Al2O6 exhibits Al6O18 puckered rings (Figure 2e) and possesses six inequivalent Sr sites which can be split in two Sr-sublattices. One of them is formed by three strontium ions six-fold coordinated by oxygen ions (Sr-O bond distances around 2.5 Å, forming irregular octahedra), and together with the two inequivalent aluminium sites, they form the framework related to perovskite materials (Figure 2c). The other three Sr ions form the second Sr sublattice. These ions are hosted within the voids between the polyhedra and also possess distorted environments, although their coordination number vary between nine, eight and seven (Sr-O bond lengths cover a wider range, between 2.43-2.93 Å). [START_REF] Alonso | Tristrontium dialuminum hexaoxide: an intricate superstructure of perovskite[END_REF] On the other hand, Sr10Al6O19

features the existence of [Al6O19] 20-unbranched hexamers rather puckered (Figure 2f), in addition to eleven inequivalent Sr sites. Two of them are six-fold coordinated by oxygen forming irregular octahedra (Sr-O bond distances range between 2.35-2.52 Å) and together with the [Al6O19] 20-hexamers form the underlying perovskite structure through corner shared SrO6 and AlO4 units. As in the case of Sr3Al2O6, the remaining strontium cations are located within the voids and feature rather irregular coordination environments (they are eight to nine coordinated, and the Sr-O bond lengths vary between 2.3-3.3 Å). [START_REF] Kahlenberg | Synthesis and crystal structure of Sr10Al6O19: a derivative of the perovskite structure type in the system SrO-Al2O3[END_REF] Overall, both compounds present two kind of Sr, the ones building up the underlying perovskite structure (i.e. six coordinated), while other Sr ions are placed within the voids exhibiting much more distorted surroundings.

wt%). [010] projections of the Sr3Al2O6 (c) and Sr10Al6O19 compounds (d). (e) View of the Al6O18 ring. characteristic of Sr3Al2O6 (f) View of the [Al6O19] 20-hexamers presented in

Sr10Al6O19. The blue tetrahedra and the green octahedra represent AlO4 and SrO6 units. Red and green balls denote oxygen and Sr atoms, respectively.

Chemical composition and fiber microstructure

In all the as-grown Sr3Al2O6 fibers prepared while varying the pulling conditions, SEM micrographs clearly reveal some bright elongated parts looking like "veins" running along the pulling direction inside a darker matrix as shown in Figure 3. The veins are depleted in Al while the Sr content is kept constant in both the veins and the matrix. Such unconventional microstructure has already been observed in Bi2ZnB2O7 fibers grown by the micro-pulling down and LHPG technique. In this latter case, the origin of this peculiar feature was attributed to the evaporation of a substantial amount of boron oxide from the melt that locally changes the chemical composition in the ternary system and leads to incongruent melting behavior. [START_REF] Assi | Towards good quality Bi2ZnB2O7 fibers grown by the micro-pulling down technique[END_REF] This further supports our previous hypothesis about the formation of Sr10Al6O19 through Al2O3 evaporation during the drawing process assisted by laser heating. Indeed, according to X-Ray spectral analysis (EDS-SEM) and elemental maps, the chemical composition of the veins and the matrix slightly differs: a Sr/Al ratio of 1.50 (perfectly matching the Sr3Al2O6 composition) and Sr/Al=1.62 are found for the matrix and the vein, respectively. The latter ratio is totally compatible with the presence of some Sr10Al6O19 secondary phase (veins) within the Sr3Al2O6 matrix, as evidenced by XRD on a crushed fiber.

Al Sr

Figure 3: EDS-SEM elemental maps of the LHPG as-grown fiber

In order to find out the effect of the pulling process over the crystal growth, i.e. crystal orientation(s), EBSD data were acquired. The analysis revealed a complex multi-scale microstructure of the LHPG fibers. While the matrix exhibits a unique crystallographic orientation, the elongated veins parallel to the pulling direction are made of several domains with different crystallographic Sr3Al2O6 orientations as shown in Figure 4a-4d. Even if the microstructures of the fibers grown from glass and crystalline seeds exhibit similar features, there exists a major proportion of veins, being more branched, for the former case (around 12±3% and 19±3 vol.%) which result in shorter fibers. Figure 4e shows the superimposition of two crystal indexations: Sr3Al2O6 in red and Sr10Al6O19 in green color. From the Image J image treatment software, the Sr10Al6O19 crystalline fraction in the fiber is estimated at 3%. This value is close from the 2.7 vol% found from the XRD Rietveld refinement. The microstructure of the ceramic bulk is depicted in Figure 4e. The ceramic is mostly made of small polycrystalline grains with an average size around 3 µm. 5). [START_REF] Black | Hydration of tricalcium aluminate (C3A) in the presence and absence of gypsum-studied by Raman spectroscopy and X-ray diffraction[END_REF] The Raman bands of the veins fit perfectly the Sr3Al2O6 cubic phase as shown in Figure 5 except for two bands. We can thus assume that these bands could be assigned to the terminal AlO4 units present in the Sr10Al6O19 composition identified as impurity by XRD and also evidenced by EBSD experiment in the "veins". As the mode at 756 cm -1 is more intense, this wavenumber range has been used for the study by principal component analysis. As it can be seen on Figure 6b the first component (PC1) is mainly affected by the intensity of the mode localised at 756 cm -1 , as the second component (PC2) (Figure 6c) is mainly affected by the frequencies change of the 740 cm -1 vibrational mode. In this way, the Raman map reconstructed on the PC1 clearly localised the Sr10Al6O19 phase within the vein as it is underlined by the color contrast of the map. PC2 points out the modification of the frequency of the main mode link to the stretching mode Al-O within the aluminium polyhedra of the Sr3Al2O6 phase. The Raman map clearly localized this local stress at the vicinity of the vein (interface vein/matrix). We can suppose that this local constraint is probably due to the fibering process which is very severe with extremely large temperature gradients at the solidliquid boundary during growth.

The interface vein/matrix was inspected by HRTEM (Figure 6). At low magnification, we see large lines oriented along the grain elongation axis and therefore potentially along the drawing axis. The orientation is always the same but we can see on the associated Fast Fourier Transform (FFT) pattern stretched diffraction spots which indicate some stacking structural defects. We can also observe the formation of rather narrow columnar domains due to the elongation of the fiber. This columnar structure with a very narrow dimension in thickness compared to the length may induce the stretched electron diffraction spots. At high resolution, the domains are slightly disoriented with each other although the general orientation remains the same. The vein is made up of a multitude of nanograins (Ø ≈ 5nm) in a little residual amorphous matrix. In HR-TEM, we can observe the cubic structure while in STEM-HAADF one can discern two types of sites. The brightest ones form large hexagons occupied only by Sr atoms (green), while the darker sites are made up of a mixture of Sr / Al (purples) atoms and is less dense. In addition, the FFT pattern corresponding to the vein shows diffraction rings which are signing the presence of crystalline nanoparticles.

The elemental compositions of the grain and veins have also been determined by STEM-EDS analyses (Figure 6). A Sr/Al ratio of 1.45 and Sr/Al=1.55 are found for the matrix and the vein, respectively, which agree with the presence of Sr3Al2O6 and Sr10Al6O19, respectively.

Finally, some precipitates enriched in Sr and Er that cross upper Sr3Al2O6 grain and the vein are observed (green circle Figure 6a) the emission spectra for two LHPG fibers grown from ceramic and glass seeds to the PL emission spectrum collected in the case of a "reference" bulk transparent polycrystalline ceramic (Cf. Figure 7a). On one hand, one can first notice the strong similarities between all the spectra, simply suggesting similar crystal field effects due to similar environments of Er 3+ .

In addition, it is worth underlining that all the emission spectra recorded for our samples compare well to the ones reported by Singh et al. for Er 3+ -doped and Er 3+ /Yb 3+ -codoped Sr3Al2O6 samples. [START_REF] Singh | Radiative properties of Er 3+ doped and Er 3+ /Yb 3+ co-doped Sr3Al2O6 phosphors: exploring the usefulness as a phosphor material[END_REF] On the other hand, one could also point out the discrepancies in the relative intensity, as well as the minor spectral differences. Indeed, the various Stark splitting are better defined in the case of the fibers compared to the "equivalent" bulk ceramic, thus clearly reflecting the superior and excellent crystalline quality of the grown materials, as already testified by the XRD data and HRTEM observations. Such moderate discrepancies probably originate from the fact that, whether Er 3+ ions populate Sr3Al2O6 or Sr10Al6O19 phases, they are still positioned at the Sr sites (Al sites being far too small to accommodate rare earth ions).

Although there exist several Sr sites (six inequivalent sites for Sr3Al2O6 and eleven for Sr10Al6O19), these environments are similar in both structures, as aforementioned. Furthermore, the PL decay curves data (see below) did not reveal any evidence for the existence of two different lifetime values which would reflect two different environments for Er 3+ ions. value equal at least to 0.997. Thus, for LHPG fibers, macroscopic PL decay measurements clearly did not reveal any evidence for an additional lifetime value that would be characteristic of Er 3+ ions localized in different environments (for instance, Er 3+ ions present within the veins, in Sr-rich precipitates and/or in the multitude of the nanograins and/or in the residual amorphous matrix -Cf. HRTEM observations).

In addition, it is noteworthy that the PL decay curve relative to the bulk transparent ceramic sample (Figure 7b) is the only one which does not correspond to a straight line. In that case, two different exponential decay functions will be necessary in order to correctly fit the experimental data (a longer lifetime value, 1 ~ 7.60 ± 0.05 ms, and another one much shorter, 2 ~ 1.75 ± 0.05 ms). The first lifetime value is very similar to the one encountered for the bulk glass sample (7.80 ± 0.05 mssee Table 1): it reflects the presence of Er 3+ ions perfectly dispersed within the ceramic grains. The second lifetime value is by far shorter, and reflects segregation/clustering effects that will occur while transforming the parent glass into a fully crystallized transparent ceramic. In optical ceramics, depending on its chemical nature, it is indeed known that segregation of the rare earth ions might occur at grain boundaries and/or that dopant concentration might vary from one grain to another. [START_REF] Zhao | Ce 3+ dopant segregation in Y3Al5O12 optical ceramics[END_REF] Thus, some moderate segregation of the Er 3+ ions is likely at play during the full congruent crystallization process, whereas such effects are not at all evidenced for LHPG fibers.

Considering that the bulk ceramic sample behaves differently than the others (two lifetimes vs a unique one), in order to be able to compare all the samples together, we have calculated the average lifetime values <  > using the following equation:

< 𝜏 > = ∫ 𝐼(𝑡). 𝑡. 𝑑𝑡 𝑡 0 ∫ 𝐼(𝑡). 𝑑𝑡 𝑡 0
All the extracted data are gathered in Table 1, and one can notice the consequent reduction (~ -20 %) of the average lifetime value concomitant with the full congruent crystallization process (from 7.80 to 6.25 ms). Finally, what must be primarily highlighted is the exceptional long fluorescence lifetime values (up to more than 13.5 ms) measured for the 4 I13/2 excited state in the case of LHPG fibers, in comparison with some literature data for equivalent or similar doping concentrations: 6.83 ms for Sr3Al2O6, [START_REF] Singh | Radiative properties of Er 3+ doped and Er 3+ /Yb 3+ co-doped Sr3Al2O6 phosphors: exploring the usefulness as a phosphor material[END_REF] 8 ms for SrGdGa3O7 single-crystal [START_REF] Piao | Optical and Judd-Ofelt spectroscopic study of Er 3+ -doped strontium gadolinium gallium garnet single-crystal[END_REF] 6.5 ms for YAG crystals, 5 just below 7 ms for Er 3+ :YAG single-crystal fibers also grown by LHPG, 4.3 ms for LiNbO3 bulk material [START_REF] Chi-Hung | Evaluation of absorption and emission cross sections of Er-doped LiNbO3 for application to integrated optic amplifiers[END_REF] and 5.75 ms for Er 3+ -doped Y3Ga5O12 nano-garnets. [START_REF] Venkatramu | Synthesis, structure and luminescence of Er 3+ -doped Y3Ga5O12 nano-garnets[END_REF] For LHPG fibers, the lifetime value was more than doubled in comparison with the average lifetime value <  > measured for the bulk ceramic sample (6.25 ms); the latter value being rather comparable to that of 6.83 ms reported in . [START_REF] Singh | Radiative properties of Er 3+ doped and Er 3+ /Yb 3+ co-doped Sr3Al2O6 phosphors: exploring the usefulness as a phosphor material[END_REF] Again, this reflects the excellent crystalline quality evidenced by HRTEM observations, and definitely constitutes serious potentialities of the grown LHPG fibers for laser applications.

Ultimately, when comparing both spectral and lifetime data, it clearly appears that the performances of our LHPG fibers are rather comparable, whatever the nature (glass or ceramic) of the seed. Nevertheless, the use of the different seed nature should be tempered. The employ of crystalline seed leads to longer fibers and better crystalline quality that slightly extent the Er 3+ emission lifetime. Nevertheless, the employ of a Sr3Al2O6 glass seed could be an attractive alternative to simplify the overall processing of the fibers, even though our experiments indicate a larger propagation of veins resulting in shorter fibers.

Conclusion

Sr3Al2O6 transparent optical fibers were successfully prepared by LHPG technique. The microstructure of the fiber is rather complex and features, in the main, a single orientated Sr3Al2O6 bulk together with polycrystalline veins with Sr10Al6O19 composition. Either glass and polycrystalline Sr3Al2O6 seeds were employed, although a large proportion of veins were observed when the glass seed was used which hinder the drawing process. The same technique turned out to be very useful for the synthesis of 1 at% Er 3+ -doped Sr3Al2O6 and the resulting fibers exhibit an intense luminescence characterized by a relatively long lifetime (13.5+/-0.1 ms) when comparing to other Er 3+ -doped materials. The transparency as well as the photoluminescent properties of Er 3+ -doped Sr3Al2O6 fiber suggest it to be an excellent candidate material for optical applications.

Figure 1 :Figure 2 (

 12 Figure 1: Photographs of the as-grown Sr3Al2O6 fiber and the LHPG process.

Figure 2 :

 2 Figure 2: (a-b) XRD patterns of the pedestal (extruded rods) and as-grown LHPG fiber.

Figure 4 :

 4 Figure 4: BSE-SEM images and corresponding EBSD maps of polished as-grown LHPG fiber

Figure 5 :Figure 6 :

 56 Figure 5: Raman spectra of the LHP fiber (vein and matrix) and the Sr3Al2O6 polycristalline

Figure 6 : 3 . 3

 633 Figure 6: (a) HRTEM and (b) STEM-EDS observation of an interface grain/vein

Figure 7 :

 7 Figure 7: (a) Normalized PL intensity relative to the 4 I13/2 -4 I15/2 electronic transition:

Table 1 .

 1 Fluorescence average lifetime values of the initial 4 I13/2 level for the four samples analyzed in this work. For the bulk ceramic sample, the two different lifetime values extracted from the use of two exponential decay functions in order to fit the experimental data.

	Fiber from	Fiber from	Bulk transparent
	ceramic seed	glass seed	ceramic

Bulk glass Fluorescence lifetime (ms) of the initial 4 I13/2 level

		6.25 ± 0.05
	13.60 ± 0.10 12.75 ± 0.10	(unique average value) 7.60 ± 0.05 / 1.75 ± 0.05	7.80 ± 0.05
		Long / Shorter lifetimes

fibre Pedestal

Acknowledgements

A.J. Fernandez-Carrion gratefully acknowledges an COST Action MP1401 grant (COST-STSM-ECOST-STSM-MP1401-050617-088753).