M G Allen
email: mark.allen@astro.unistra.fr

P Fernique

T Boch

Daniel Durand

A Oberto

B Merín

F Stoehr

F Genova

F.-X Pineau

J Salgado

B Merin

An Hierarchical Approach to Big Data

published or not. The documents may come

HiPS

The Hierarchical Progressive Survey (HiPS) is a HEALPix1 (Górski et al. 2005) based scheme that enables a multi-resolution approach to astronomy data from the individual pixels up to the whole sky (Fernique et al. 2015). HiPS was initially developed at the Centre de données astronomiques de Strasbourg (CDS2 ) to enable the visualisation of all sky imaging survey data in Aladin (Fernique et al. 2010). The initial developments have been extended to astronomical source catalogue data, and also to 3-dimensional data cubes.

The cooperative use of HiPS by a number of data providers has led to a network of distributed HiPS nodes for sharing and mirroring data sets. We attribute the rapidly growing implementation of this system to the convergence of a number of technical and organisation factors that match the needs of the community. In this article we identify the approaches that have been taken to make this scheme a practical solution for managing large volumes of heterogeneous data, and we consider its applicability to large future data sets.

Design and implementation considerations

The underlying idea of an hierarchical approach to managing and visualising astronomical survey data is that the data can be accessed at the scale and resolution required for a given purpose. With a hierarchical system one can access just the data that is necessary. This has been very successful for visualisation of large surveys in tools such as Google sky3 , Microsoft World Wide Telescope4 (WWT), and Aladin5 . The HiPS scheme builds on the concepts for hierarchical visualisation of survey data with the aim of creating a scientifically robust system for hierarchical use of different data types (catalogues and cubes), and to enable easy and practical use of these data in an interoperable way. HiPS is the result of a long term effort to make data useable and interoperable. The design and implementation considerations that have proven to be successful in the development of HiPS are that: the scientific properties of the original data must be preserved; the system must be simple to use; and that the client tools for use of the system should be independent.

Preserving the scientific properties of the data has been a key point for going beyond visualisation of the data, and for addressing the requirements for scientific use. The choice of HEALPix as the tessellation for HiPS provides a strong scientific basis because of the existing software and use in the community. Also because of its scientific properties such as equal area pixels and the direct translation of indices into astronomy coordinates, as well as other properties (see - [START_REF] Reinecke | [END_REF]). HiPS does of course involve the resampling of original pixel data onto a HEALPix grid, but in order to preserve the original dynamic range of the data values HiPS tiles can be encoded as FITS files (as well as JPEG and PNG). Also, HiPS has an in-built mechanism for the direct referencing of the original data files which contribute to a given HiPS tile, providing important information on the provenance of the data values.

The development of HiPS has involved making choices on the balance of simplicity and complexity in different parts of the system. The structure and encoding of the HiPS data was chosen to be very simple, and to employ standard and pervasive technology. HiPS tiles are stored as simple files in a file system directory structure that is organised into 'HiPS orders' that represent HEALPix maps at different resolutions. This permits direct access to the necessary files for a given region of the sky at the required resolution. The use of simple files and directories (rather than databases) allows HiPS datasets to be published by simply moving the files to a http server. This choice has proved to be extremely useful and well matched to the needs of astronomers and data centre staff as this level of implementation can be performed very easily. An astronomer can create, test and publish a HiPS without need for database expertise, and data centres can test and evaluate use of HiPS with minimal investment or risk. This level of simplicity is possible because file access is sufficiently fast, and because the indexing is implicit in the definition of the HiPS hierarchical directory and file structure.

The HiPS structure has also proved to be flexible enough to expand its use to other data types based on the concept that a HiPS tile is a generic 'container' (Fig 1) which may contain any kind of information related to the relevant sky area and resolution level defined by the tile. In this way, the tiles may contain not just images, but data cubes, or catalogues or related provenance links as described in Fernique et al. (2015).

While the HiPS structure itself is very simple, the methods required to generate HiPS data sets are complex. The initial mapping of an image survey into a HEALPix grid can require careful treatment of the mosaicking and background level characteristics of the data. It was a design decision that this complexity be built into the HiPS generation tools. This gives the data providers the flexibility to define and control the way their data is mapped onto a HiPS structure, and provides a way for the data providers to decide the specifics of how the scientific properties of the data are preserved. Significant effort has been put into the development of the hipsgen tool, in particular to improve the speed of computing the HiPS tiles from the original data.

There are a number of tools that are available for the use of HiPS. The Aladin desktop java application is a full featured application that has been used as the primary test platform for HiPS. Aladin provides all sky visualisation of HiPS, supporting multiple views with synchronised zoom and pan, and also transparency overlays and interoperability with other online astronomy services. To enable widespread use and access to the tools for generating HiPS, the hipsgen capabilities are built into Aladin itself, as well as being available as command line programs6 .

Aladin Lite is a simplified HiPS JavaScript visualiser for web browsers. It is designed to be embedded on a third-party web pages [START_REF] Boch | [END_REF]) and it comes with an application program interface (API) that provides the necessary controls to customise it to different needs. Aladin Lite has been particularly important for the implementation of HiPS [START_REF] Boch | ADASS XXV[END_REF]. Examples include 'ESA Sky' [START_REF] Merin | ADASS XXV[END_REF], the Spitzer GLIMPSE 360 web page 7 and implementations at CADE 8 to preview HEALPix maps of various surveys. Aladin Lite requires only small snipets of code to be embedded, and collaboration with implementors shows that the level of customisation available matches their needs at the right level.

HiPS for Big Data

There are currently ∼250 HiPS sets available from the CDS and other data providers (see the HiPS directory 9 ). This diverse set of HiPS covers a wide range of wavebands, scales and resolution. The recently up-dated HiPS for the HST images available from CADC shows that an entire archive of pointed observations can be combined and used as an image survey. Recent tests with ALMA data cubes show that the calculation and usability of individual 5TB HiPS data sets is feasible.

Since HiPS uses individual files, its scalability for Big Data mainly concerns the issue of scalability and management of file systems that contain many small files, and we expect that generic Big Data solutions for large file systems will be applicable to 'Big HiPS'. Looking ahead, this scalability provides a way to plan for some aspects of the visualisation and data set exploration challenges that will come with LSST and SKA. A HiPS of a hypothetical LSST data subset covering 18000 square degrees of the sky at 0.3 arc-second resolution, each 3 days for 3 years would lead to a ∼5 Petabyte HiPS cube in FITS format, or 256 TB in JPEG format. HiPS access for visualisation of such volumes would be feasible today.

Figure 1 .

 1 Figure 1. HiPS multi-resolution tile structure based on HEALPix

http://healpix.sourceforge.net

http://cds.unistra.fr

http://www.google.com/sky

http://www.worldwidetelescope.org

http://aladin.unistra.fr

e.g. java -Xmx16000m -jar Aladin.jar -hipsgen in=Fits_directory

Acknowledgments. MA aknowledges support from the Astronomy ESFRI and Research Infrastructure Cluster -ASTERICS project, funded by the European Commission under the Horizon 2020 Programme (GA 653477).


Laure Caracciolo

Lénaïc Madec
email: lenaic.madec@univ-pau.fr

Emmanuel Petit

Vincent Gabaudan

Dany Carlier

Laurence Croguennec

Hervé Martinez

Electrochemical redox processes involved in carbon-coated KVPO4F for high voltage Kion batteries revealed by XPS analysis

K-ion batteries need efficient positive electrode materials with open structural frameworks to accommodate the large ionic radius of K + . In that direction, polyanionic compounds are of great interest. Among them, KVPO4F is the most studied one. Its electrochemical redox processes still remain unclear, especially at high voltage. To tackle this issue, X-ray photoelectron spectroscopy was used to reveal the electrochemical redox processes of KxVPO4F. First, a carbon coating was performed and allowed increasing the overall electrochemical performance while mitigating electrolyte degradation at high voltage. Then, XPS analysis showed a high reversibility of the redox processes although the K + extraction (and insertion) from x=0.5 to 0 was hindered, possibly by structural constraints while electrolyte degradation occurred mostly above 4.5 V.

Introduction

The development of new electrochemical energy storage systems as alternatives to the Li-ion technology remains a key challenge due to the growth of energy storage demand.

Alternative batteries based on abundant and thus low cost raw materials are required.

Na-ion and K-ion batteries have attracted increasing interest during the last decade.

Thanks to the low potential of the K + /K redox couple versus the other alkali metals in non-aqueous solvents, high voltage K-ion batteries (KIBs) could be developed. 1 Moreover, K + has the lowest Lewis acidity and desolvation energy compared to Na + and Li + , 2 and thus higher ionic conductivity and faster electrode/electrolyte interface diffusion kinetics that would allow developing high power KIBs.

The high reactivity of K metal requires considering results obtained in half-cells with caution, especially the interaction between the electrolyte and the electrode surface. Indeed, migration of electrolyte degradation products can occur from the K metal to the working electrode. 3 Among negative electrode materials, 4 graphite is of great interest as it reversibly intercalates K + to form KC8. 5,6,7 Promising electrochemical performance is obtained, with a reversible capacity of ∼250 mAh.g -1 up to hundreds of cycles and even for high practical loading (> 20 mg/cm 2 ). 7,8,9,10,11,12,13 Regarding the safety issue, the average discharge potential of ∼0.3 V vs. K + /K makes graphite even more attractive for the application.

In fact, the major drawback of K-ion batteries is the need of positive electrode materials with open structural frameworks to accommodate the large ionic radius of K + (1.38 Å) compared to Na + (1.02 Å) and Li + (0.76 Å), 1,14 which is detrimental to the volumetric energy density. 15 Polyanionic compounds are of great interest as they are suitable host structures for the reversible K + intercalation as well as to deliver high voltage. 16 For instance, KVPO4F (KVPF) showed reversible capacity up to ∼105 mAh.g -1 with an average discharge potential of ∼4.3 V vs. K + /K and excellent rate performance. 17,18,19,20,21,22 In a full-cell configuration using graphite, the energy density has even been evaluated at 300 Wh.kg -1 , suggesting a promising material for further fullcell studies. 23 Regarding its electrochemical behaviour as a positive electrode material, KVPF shows 4 distinct potential domains in both charge and discharge, corresponding to successive phase transitions with an interesting limited lattice volume change of ∼6.5 %, 19 and the formation of intermediate phases KxVPO4F at x=0.75, x=0.625 and x=0.5.

The reaction occurring at high voltage, between the composition K0.5VPO4F and VPO4F still remains unclear, however, with simultaneous active material redox reaction and large electrolyte degradation. It may explain why the theoretical capacity of KVPF (131 mAh.g -1 ) was never reached so far. Therefore, it is of great interest to better understand the electrochemical processes involved in KVPF positive electrode materials in KIBs in order to improve its performance.

The present work proposes to investigate the electrochemical redox processes involved in KVPF upon cycling in K metal batteries using X-ray photoelectron spectroscopy (XPS). The impact of a carbon-coating on the electrochemical performance of KVPF was first evaluated, regarding its ability to improve the overall performance and mitigate the electrolyte degradation.

Experimental

KVPO4F and carbon coated KVPO4F (KVPO4F-C) were prepared by a two-step reaction. First, VPO4 was obtained by mixing V2O5 (from Sigma-Aldrich, 99.6%) and NH4H2PO4 (from Sigma-Aldrich, 99.99%) by high energy ball-milling (SPEX 8000M mixer/mill at 1425 rpm) for 1hr followed by a thermal treatment at 800°C for 7h under Ar/H2 (5%). Second, stoichiometric amounts of KF.2H2O (from Alfa Aesar, 98.5%, dried overnight under vacuum at 250°C and stored in a glovebox under Ar) and VPO4 were mixed without (KVPF) or with (KVPF-C) the addition of 10 wt.% sucrose (from Sigma-Aldrich, ACS reagent). The mixture was obtained by SPEX ball-milling for 1hr, pelletized and thermal treated at 650°C for 8h under Ar. Rietveld refinement of the Xray diffraction patterns was performed using the FullProf software and considering an orthorhombic unit cell described in the Pna21 space group. 24 It confirmed the nature of the grey (KVPF) and black (KVPF-C) powders obtained (Fig. S1). The smaller cell parameters and especially volume cell obtained for KVPO4F-C versus the KVPO4F phase reported and characterized in-depth by Fedotov et al. 17 suggest a partial O 2- substitution for F -and thus the formation of KVPO4F1-yOy. y was estimated to 0.1 from the study of the solid solution KVPO4F -KVPO4O (Table S1, and from 31 P MAS NMR data not reported here). Note that the stoichiometry in K/V/P was determined by chemical analyses (Fig. S1). In addition, thermogravimetric analyses and scanning electron microscopy images confirmed the formation of a carbon coating (2.5 wt.% of C) for KVPF-C compared to KVPF (Fig. S2). Electrodes were prepared by mixing KVPF or KVPF-C, carbon black (Super P, BET=62 m 2 /g, Alfa Aesar) and poly(vinylidene fluoride) according to a 70:25:5 weight ratio, in N-methyl-2-pyrrolidone (32:68 solvent:material weight ratio), by ball-milling for 1h. The obtained slurry was cast on an aluminium current collector, dried for 24h under Ar, and finally electrodes of 1.5 mgKVPO 4 F/cm² were punch out and dried under vacuum at 80°C for 12h. 2032 cointype cells (316L stainless steel) were assembled under Ar using a Whatman (grade GF/D) and a polypropylene membrane (Celgard) as separators, K metal (99.95%, Alfa Aesar) as negative electrode and 100 μL of 0.8M KPF6 in a mixture of EC:DEC by volume as electrolyte (KPF6, Sigma-Aldrich, ⩾99% ; EC, Sigma-Aldrich, anhydrous, 99% ; DEC, Sigma-Aldrich, anhydrous, ⩾99%). All potentials reported in the text are referred to K + /K. Note that K metal polarization in symmetrical K//K cell was found negligible during open circuit voltage (< 6 mV, Fig. S3) compared to the literature. 25 During plating/stripping at 0.04 mA/cm 2 (corresponding to the current used to cycle K//KVPF cells), however, a polarization of 0.1 V was measured. Cycling tests were performed at C/5 (26 mA.g -1 , i.e. the exchange of 1K + in 5 hours) and 20°C using a VMP3 (Biologic, SAS France). Two formation cycles were first performed between 3.5-5 V vs. K + /K to minimize electrolyte degradation impact (Fig. S4 and Fig. 3), as discussed later. KVPF-C cells were then stopped at different potentials during the 3 rd cycle and electrodes were recovered from the cells and washed twice with 1 mL of DEC during 30s. XPS analysis was performed using an Escalab 250 Xi spectrometer, with a monochromatized Al Kα radiation (hν = 1486.6 eV). Electrodes were placed on a sample holder using uPVC insulation tape (3M part number 655), and transferred to Ar filled glove box connected to the spectrometer. Using the standard charge compensation mode, core spectra were recorded with a 0.15 eV step and a constant 20 eV pass energy. Using CasaXPS software, the binding energy scale was calibrated from the KVPF oxygen peak at 530.0 eV. It is reminded that XPS probes about 5 to 10 nm depth of the sample surface, and here between 6-7 nm for the vanadium, i.e. not only the extreme surface atomic layers (0.5 -1 nm) but also the deeper ones, which means a representative part of the bulk is also probed (Fig. S5). Moreover, at the extreme surface, the presence of defects may lead to vanadium with an oxidation state higher or lower than in the bulk. V 2p3/2 core spectra of KVPF electrodes were thus fitted using peak positions and line shapes constraints corresponding to different oxidation states as determined from reference samples with V n+ in a PO4 environment (Fig. S5 andTable S2) in agreement with the literature. 26,27,28,29,30 This thorough procedure allows reliably following the evolution of the mean oxidation state upon cycling, especially considering that the mean oxidation state obtained by XPS (3.1) for the pristine KVPF-C material is in good agreement with the XRD analysis (KVPO4F1-yOy with y=0.1).

Results and discussion

The impact of the carbon coating was evaluated regarding its ability to optimize the electrochemical performance of the material in KIBs and to mitigate the KVPF active material surface reactivity against the electrolyte. Overall, both KVPF and KVPF-C showed a similar charge/discharge profile at C/5 with 4 potential domains (Fig. S4), in very good agreement with the literature. 19 However KVPF-C showed a lower average polarization of ∼230 mV compared to ∼300 mV for KVPF (Fig. S4) and a better capacity retention after 70 cycles compared to KVPF (76 % vs 62 %, (Fig. S6). For both systems, a significant high-potential hysteresis is observed, suggesting that the main capacity loss would be related to the high-potential domain. Since KVPF and KVPF-C have close particles size (Fig. S2), the carbon coating was shown -as expected -to promote better transport properties within the positive electrode, lower polarization of the cell, smaller electrolyte degradation at high voltage and thus optimized performances. XPS analysis confirmed a significantly lower solid electrolyte interphase coverage at the KVPF-C surface compared to KVFP after the first cycle (Table S3). Note that the lower electrolyte degradation is explained by a passivation effect of the carbon coating at the KVPF surface. Based on these results, KVPF-C was therefore selected to study the electrochemical redox processes involved in KVPO4F upon cycling. Fig. 1a shows the typical 3 rd charge/discharge cycle obtained for K//KVPF-C cells and indicates at which potentials the prepared cells were stopped to perform the analysis of the redox processes involved in KVPF-C. Note that charge/discharge profiles obtained for all the cells showed very close polarization and capacities (Fig. S7), highlighting the great reproducibility of the electrochemical results so that reliable XPS analysis is expected on materials recovered ex situ from this series of cells. Fig. 1b shows the V 2p3/2 XPS core spectra of KxV n+ PO4F-C electrodes as recovered from K//KVPF-C cells during the third cycle at C/5, as function of the K content. For more information, the full V 2p core spectra can be found in Fig. S8. XPS analysis revealed that the pristine KVPF-C showed the presence of V 2+ , V 3+ and V 4+ (Fig. 1b) although only V 3+ and V 4+ were expected based on the KVPO4F0.9O0.1-C formula. Note that V 2+ was also observed for the VPO4 precursor (Fig. S5). However, considering the very low amount (< 1wt.%) of VPO4 impurity in the KVPF-C material (Fig. S1), the presence of V 2+ is likely to originate from the KVPF-C extreme surface. This was supported by the decrease of the V 2+ contribution during oxidation as discussed thereafter.

After charge of the cell to 4.10 V, i.e. during the 1 st oxidative potential domain, the peak associated with V 4+ increased while the V 3+ and V 2+ ones decreased (respectively +7%, -4% and -3%). At almost the end of the 2 nd potential domain at 4.30 V (i.e. at the composition of about K0.625VPO4F), the peak associated with V 4+ increased while the peaks associated with V 3+ and V 2+ decreased (respectively +15%, -12% and -3%). At the end of the 3 rd potential domain at 4.46 V (i.e. for approx. K0.5VPO4F), a new peak, attributed to V 5+ appeared (+13%) while the V 4+ , V 3+ and V 2+ peaks decreased (-11%, -1% and -1%, respectively). This phenomenon further continued after the 4 th potential domain at 5.00 V, so that the V 5+ , V 4+ , V 3+ and V 2+ peaks represent 20%, 50%, 23% and 7%, respectively. Note that despite the remaining V 3+ and V 2+ even at the end of charge, the corresponding mean oxidation states were in good agreement with the theoretical/electrochemical ones, as discussed thereafter. Note that remaining of V 3+ and V 2+ at the end of charge should arise from specific vanadium environment at the KVPF surface. During discharge, opposite phenomena were observed (Fig. 1b), indicating a high reversibility of the redox processes involved in KVPF-C upon cycling in the potential window 3.5-5V vs. K + /K. The evolution of the vanadium mean oxidation state as the K content changed during the 3 rd cycle for KVPO4F0.9-O0.1C and determined from both the charge/discharge capacity (DOelec) and XPS analysis (DOxps) is reported in Fig. 2.

Until the end of the 3 rd potential domain, i.e. at 4.46 V, a very good agreement was observed between the vanadium mean oxidation states determined from the electrochemistry and from XPS analysis. On the contrary, at the end of the 4 th potential domain, i.e. at 5.00 V, DOelec was +3.9 while DOxps was only +3.8. This discrepancy indicates that about 25% of the electrons exchanged in the last potential domain are not related to K + deintercalation from KVPF-C but to electrolyte degradation. This is in agreement with the high irreversible capacity observed for the 4 th potential domain between the charge and the next discharge (Fig. 1a). Moreover, the evolution of the capacities associated with the different potential domains for the first 20 cycles showed that electrolyte degradation is mainly associated with the 4 th potential domain as shown in Fig. 3 by the large difference in height between the full and hatched black bars.

Nevertheless, during the first two cycles, irreversibility was observed throughout the reaction, the latter being much more pronounced at high voltage (about 60%, Fig. 3).

During the next cycles, a stabilization was observed with an excellent efficiency below 4.6 V (same height for the full and hatched blue and green bars), and thus a capacity loss almost exclusively associated with the high voltage domain (about 90%) and the composition range K0.5VPO4F-C -VPO4F-C. This highlights the relevance of using the 3 rd cycle to perform the redox processes study instead of the 1 st cycle. During discharge, DOxps decreased back to +3.1 (Fig. 2), indicating a complete reversibility of the redox processes. This result fully supports the limitation of the participation of the electrolyte degradation to high voltage, only after the 2 nd cycle as shown in Fig. 3.

Importantly, only 0.7 K + were extracted from KVPO4F0.9-O0.1-C (Fig. 2) instead of 0.9 theoretically. Especially, the capacity associated with the 1 st , 2 nd and 3 rd potential domains nearly reached the theoretical ones and remained almost constant from cycles 3 to 20 (Fig. 3). On the contrary, the 4 th potential domain capacity continuously decreased to about half the theoretical one from cycle 3 to 20 (Fig. 3). Interestingly, no polarization increase (due to the accumulation of electrolyte degradation products) was observed (Fig. S9). Thus, the 4 th potential domain capacity decrease is not explained by the cut-off potential being reached before the end of the domain. Instead, these results likely originate from structural constraints hindering the extraction (and insertion) of K + from x=0.5 to x=0 as proposed previously. 19,31 To evaluate the possibility to further extract the remaining K + , a cell was hold at 5 V for 18h during the 3 rd cycle and the corresponding DOxps was +3.9 (Fig. S10), in agreement with the evolution of the cell parameters determined by XRD (not shown). This result highlights that despite the electrolyte degradation process, it is still possible to extract the potassium but the process is much slower, more likely due to a structural constraint such as cell volume reduction. This structural constraint could increase the K + migration barrier, inducing a decrease in the K + diffusion coefficient, 31 slowing the global extraction process.

Conclusion

In conclusion, the carbon coating of the KVPF material was found beneficial to the overall electrochemical performance while mitigating electrolyte degradation at high voltage. Then, by comparing the average oxidation state derived from electrochemical and XPS analysis, the electrochemical redox processes of KxVPO4.1F0.9-C were revealed.

Interestingly, a high reversibility was found although the K + extraction (and insertion) from x=0.5 to x=0 was hindered, possibly by structural constraints. Moreover, the severe electrolyte degradation issue, observed mostly above 4.5 V, is currently under further investigation, as well as possible KVPF dissolution in order to successfully mitigate these irreversible degradation reactions using electrolyte additives and enable highly efficient KVPF based full K-ion cells. 3 Capacity (mAh.g -1 ) related to the 1 st , 2 nd , 3 rd and 4 th redox processes as obtained from the derivative of the galvanostatic charge/discharge curves. Note that the 2 nd and 3 rd potential domainss were combined as they cannot be perfectly dissociated after few cycles (not shown).

Corresponding capacity loss (%) between oxidation and reduction is also reported by the curves with the red points. Theoretical capacity of the different redox processes based on the KVPO4F0.9-O0.1C formula (118 mAh.g -1 ) are also reported for comparison.

Fig. 1 a

 1 Fig. 1 a) Typical 3 rd charge/discharge cycle (at C/5, 26 mA.g -1 ) obtained for K//KVPF-C cells indicating at which potentials cells were stopped for the analysis of the redox processes involved in KVPF-C. b) Vanadium 2p3/2 XPS core spectra of KxV n+ PO4F0.9O0.1-C electrodes as recovered from K//KVPF-C cells during the 3 rd cycle, as function of the potential.

Fig. 2 .

 2 Fig. 2. Vanadium mean oxidation state (DO) of KVPO4F0.9-O0.1-C, as the K content changed during the 3 rd cycle, determined from the charge/discharge capacity on one side and from the analysis of the XPS spectra on the other side. Grey lines represent the theoretical DO based on the KVPO4F0.9-O0.1-C formula.

Fig.

  Fig.3Capacity (mAh.g -1 ) related to the 1 st , 2 nd , 3 rd and 4 th redox processes as obtained from

Acknowledgments

The authors thank the financial support of E2S-UPPA, Région Nouvelle Aquitaine and the French National Research Agency (STORE-EX Labex Project ANR-10-LABX-76-01 and TROPIC project ANR-19-CE05-0026-01). The authors thank Cathy DENAGE, Etienne DURAND and Eric LEBRAUD (ICMCB) for their help with SEM, TGA-FTIR and XRD analyses, respectively. Authors also thank Joël GAUBICHER (IMN-Nantes) for fruitful discussions about redox processes in KVPF.

Conflicts of interest

There are no conflicts of interest to declare.


Leila Ben Amor

Besma Belgacem

Jean-Sébastien Filhol

Marie-Liesse Doublet

Mouna Ben

Rached Ben Hassen

New p-type Al-substituted SrSnO 3 perovskite for TCO applications?

Novel p-type SrSn 1-x Al x O 3 (x=0, 0.2, 0.5) perovskites are presented as potential candidates for electro-optical applications. A combined experimental and theoretical study reveals that chemical substitutions can be used as a lever to stabilize oxygen holes in the valence band.

Transparent conductive oxides (TCO) are an original class of materials combining two contrasting properties in one single compound that is, electrical conductivity and optical transparency in visible spectrum. TCOs are widely used in electro-optical devices such as flat-panels, organic light emitting diodes, photovoltaic devices, and plasma displays.

1-6 However, standard commercial TCOs are limited to post-transition metal oxides, e.g. ZnO, In 2 O 3 and SnO and derivatives.

 due to the strongly localized states of the acceptor, i.e. oxygen, which are at the origin of large hole effective mass (poor electronic conductivity). As a consequence, the probability is high that the newly created shallow acceptors are counterbalanced by the formation of native defects such as anion vacancies or cation interstice.

12 Up to now, n-type tindoped indium oxide (ITO) shows the best performance.

1,13

Nevertheless and despite the remarkable properties of ITO, the relative high cost of In and its increasing scarcity appeals for alternative elements or materials. 14 ABO 3 perovskites are of great interest thanks to the use of mixed A and B cations capable of generating versatile structures and properties with high chemical stability. SrSnO 3 (SSO) is one of the most promising candidate due to its high visible light transmittance, high thermal stability and reasonable cost. 15 Pure SSO is a charge transfer insulator crystallizing in the orthorhombic GdFeO 3 -type structure with lattice parameters a = 5.7113 Å, b = 8.0647 Å and c = 5.7042 Å. In this tri-dimensional structure, SnO 6 octahedra are corner-shared and Sr lies in the available vacant interstices. 16,17 To meet the industrial specifications of efficient TCOs, SSO must display high carrier concentration and mobility [18][19][20][START_REF] Woodward | The Electronic Structure of Metal Oxides[END_REF] which requires its doping. Attempts to develop n-type SSO candidates that could rival the electro-optical properties of ITO were intensively described and interesting performance were reported. [22][23][24][25][26] In contrast, p-type doping is still lacking for this class of materials. 27-29 Recently, Hautier et al. showed that the large difference in mobility between n-type and p-type materials is not inevitable and could be overcome by the investigation of alternative chemistries. Based on this statement and considering the high industrial stakes associated with the discovery of new p-type TCOs, we used chemical substitution to check the ability of SSO to behave as a performant p-type TCO. This concept was introduced for the first time by Hosono and coworkers [START_REF] Hosono | Proc. Sympodium O Thin Film Chalcogenide Photovolt. Mater. EMRS 2006 Conf[END_REF] and is known as the "chemical modulation of the valence band" (CMVB). Based on the hypothesis that partial substitution of Sn(IV) for the smaller trivalent cation Al(III) can enhance electron delocalization (through a significant volume contraction of the perovskite structure) while concomitantly generate stable holes in the valence band, various Sn/Al-substituted SSO phases were synthesized. Using X-ray diffraction (XRD) and Scanning Electron Microscopy (SEM) equipped with Energy Dispersive X-ray Spectroscopy (EDX), we here demonstrate the solubility of the Al element in the SSO perovskite structure for SrSn 1-x Al x O 3 (SSAO) compositions x = 0 to 0.5. UV-Vis-NIR spectroscopy is then used to show that the absorption profile of SSO is not altered by the chemical Sn/Al substitution while the lowenergy absorption band (intra-band transitions) is enhanced. A thermodynamic study based on first-principles DFT (Density Functional Theory) calculations is then performed to determine the type of charge carriers in SSO and SSAOs (p-type) and combined with chemical bond analyses to rationalize the impact of Sn/Al substitution on the p-type performance of SSAOs. SrSn 1-x Al x O 3 (x= 0, 0.2, 0.5) perovskites were successfully obtained through solid state reaction. Details of the synthesis steps are given in ESI part. The impact of the Sn(IV)/Al(III) substitution ratio on the macroscopic properties of SrSn 1-x Al x O 3 was investigated through complementary XRD and SEM-EDX. The structural and morphological analyses of the synthesized samples are presented in Fig. 1 andS1. From the XRD patterns of SSO and SrSn 1-x Al x O 3 (SSAO) the orthorhombic Pnma structure (JCPDS 10-6354) is confirmed. A slight shift towards higher angles is observed in agreement with the smaller Shannon ionic radius of Al 3+ (0.54 Å) compared to Sn 4+ (0.69 Å). For all compositions, the Goldschmidt tolerance (t) and distortion octahedral () factors remain in the range [0.8-1.1]

30 and [0.41-0.73], respectively. [START_REF] Wang | Functional and Smart Materials[END_REF][START_REF] Zhu | [END_REF] which confirms the phase stability. Note that a small amount of SrCO 3 impurity was detected and attributed to the potential reaction of Sr defects with atmospheric oxygen and carbon dioxyde. 33 This is confirmed by DFT calculations showing that the formation of SrCO 3 from SSO is favoured for all compositions (see Table S1). Interestingly, the content of SrCO 3 impurities slightly decreases with the increase of Sn/Al ratio (xAl Sn ). This suggests that Sr vacancies are less favoured in SSAO as xAl Sn increases, as confirmed by DFT calculations (see Table S1). For xAl Sn = 0.5, extra reflections appear in the XRD patterns that could be assigned to either a superstructure, or the formation of a cubic-type structure as previously reported by Waerenborgh and coworkers for xAl Fe =0.5 substitution in SrFe x Al 1-x O 3 . 34 Structural predictions carried out by means of a basing-hopping procedure clearly identify cubic SSAOs as thermodynamically favoured over orthorhombic polymorph when xAl Sn reaches x=0.5. This could arise from an Al ordering occurring in the structure (see Fig. S2). The magnified scanning electron microscopy (SEM) images and energy dispersive x-ray (EDX) cartographies given in supplementary information (Fig. S1) show the formation of micron-sized agglomerates of raw elements Sr, Sn, Al and O. An additional carbon peak is also observed which originates from the sampling method that uses graphite film deposit. Interestingly, when measurements are performed on pellet, the carbon signature drastically decreases and SrCO 3 impurities remain as trace (see Fig. S1d). UV-Vis-NIR Optical reflectance was further measured for each sample up to 2700 nm (Fig. 2a). All spectra are similar and show two distinct zones: one strong optical absorption zone in the range 270-400 nm, which is attributed to the transition from the VBM (predominant O(2p) states) to the CBM (predominant Sn(5s,p)). Wide and less intense bands are then observed around 500 nm that can be assigned to carbonate-type impurities (known to absorb in that region 35 ). The amplitude of these bands decreases with the increase of xAl Sn , consistently with the progressive decrease of the amount of SrCO 3 impurity. Interestingly, additional peaks are observed in the 1400-2700 nm NIR region which were not reported in the literature, even for the un-doped SSO material. The shape of these peaks is very similar for SSO and SSAO suggesting equivalent intra-band electronic transitions and equivalent type of charge carriers for all compositions. Noteworthy, the width and intensity of these transitions is directly correlated to the Sn/Al substitution ratio. This absorption continuum indicates interesting conducting properties, possibly associated with shallow acceptor states (i.e. stable oxygen-holes) in the valence band, both for unsubstituted SSO and substituted SSAOs. To determine which charge carriers occur in SSAOs (n-or p-type), electronic conductivity and Seebeck (or Hall-effect) measurements are necessary. However, such measurements require high-temperature sintering techniques to get high-density samples. This procedure is known to induce undesired O 2 release that can switch transport properties from p-to n-type, therefore leading to unreliable results. 36 Indeed, the main impediment to p-doping in oxides has been identified as the formation of compensating intrinsic defects (hole killers) such as oxygen vacancies. 12,37,38 The occurrence of stable oxygen vacancies (V O ) is then not desired as they may kill the shallow acceptor states created by Sr-defects or Al Sn substitution. In contrast, native defects of cations such as Sr (V Sr ) and Sn (V Sn ) are beneficial for p-type behavior. It is therefore crucial to determine the nature of the defects in the SSO and SSAO structures in order to properly assign the low-energy transitions observed in their absorption spectra. For this propose, a thermodynamic DFT study was conducted for the substituted SSAO phases to assess the stability of the oxygen network towards O 2 gas release, with and without the presence of Sr vacancies.

𝑆𝑟 1-𝑦 𝑆𝑛 1-𝑥 𝐴𝑙 𝑥 𝑂 3 → 𝑆𝑟 1-𝑦 𝑆𝑛 1-𝑥 𝐴𝑙 𝑥 𝑂 3-𝛿 + ( 𝛿 2 ) 𝑂 2 (1)
According to reaction (1), a direct relation can be established between 𝑦, 𝑥 and  to identify the doping-type of the material. A ptype behavior (resp. n-type) is expected for 𝑦 + 𝑥 2 > 𝛿 (resp. < 𝛿) i.e. when the charge compensation due to Sr vacancies and Sn/Al substitutions remains greater (resp. lower) than the content of released O 2 . The frontier between both regimes is illustrated by a dashed line in Fig. 3 and corresponds to a number of holes per O of 0.087. Remarkably, the reaction enthalpy remains positive up to 0.13 holes per oxygen, which suggests that p-type SSAOs could be stable in a narrow domain of O-holes from 0.087 to 0.13, with no oxygen release. Above this critical value, O 2 is released but the system remains p-type (𝛿 < 𝑦 + U(1/2,0,1/2) and (0,0,0) consistent with previous reports. 41,42 Interestingly, SSAO displays a direct band gap around  showing that the symmetry lowering due to Sn/Al substitution does affect the general shape of the band structure. The computed band gaps are found to be 4.3 and 4.6 eV, for SSO and SSAO with xAl Sn = 0.25, respectively, in good agreement with recent calculations using the hybrid HSE (Heyd-Scuseria-Ernzerhof) functional, 43 and in fairly good accordance with the experimental blue-shift observed in Fig. 2b. The shallow states observed at the VBM of SSAO increase with the Sn(IV)/Al(III) substitution ratio. According to the projected Density of States (DOS) these states are assigned to the oxygen atoms that are directly connected to Al (Fig. S4). This is fully consistent with the less efficient electrostatic stabilization of O . 44 The comparison of the electronic band structures of Figs. 4a andb highlights a perturbation of the electronic states of SSAO compared to SSO and a substantial increase of the valence bandwidth. To quantify the VBM dispersion and transport properties, anisotropic hole effective masses were evaluated for SSAO by measuring the band curvature around , Y and Z , as well as along the most interesting -X (½, 0, 0), Y(0, ½, 0)-M(½,½,0) and Z(0, 0, ½ ) -R(0, ½, ½) directions of the Brillouin zone (See Fig. S5). Results give 0.62m e , 0.43m e and 0.67m e for SSAO and 0.80m e , 0.50m e and 1.76m e for SSO, respectively, showing that SSAO (xAl Sn =0.25) displays lower effective masses than SSO for a given direction. Heavier effective mass is obtained on U and along U- direction (1.34m e for SSAO and 4.77me for SSO). This can be fully rationalized via chemical bond analyses through the computation of Fukui functions (f -) and electron localization functions (ELF). As detailed in SI, these topological tools of the electron density provide direct proves of the VBM states (Fukui) and real-space Lewis-like structures (ELF). The Fukui functions of SSO and SSAO reveal that the VBM states have an oxygen character with, however, very different spatial distributions for both systems. While SSO shows O(2p) orbitals perpendicular to the (a,c) plane, a more spherical density is clearly visible around the oxygens for SSAO, indicative of an enhanced Sn-O bond covalency (Figs. 4c-d).

The ELF plots of Fig. 4g-h show that two oxygen lone-pairs pointing towards the Sr 2+ cations (two yellow volumes) for both SSO and SSAO. They correspond to the O(2s) and O(2p) lone-pairs of oxygens that are not involved in covalent bonds with Sn or Al and are consistent with the oxygen Lewis structure that is expected in such local environment (see Fig. 4g). Interestingly, another lobe is visible along the Al-O bonds but close to O, which illustrates a loss of Al-O interactions in the lattice, fully consistent with the less electronegative character of Al with respect the Sn (Fig. 4h). Owing to the smaller radius of Al compared to Sn, the Al/Sn substitution induces a significant contraction of the Sn-O bonds, from 2.08 to 2.04 Å and a substantial O-Sn-O angle increase from 155° to 175° which both contribute to the increase of the valence band width (Fig. 4d). These local distortions impact the cell parameters of SSAO which notably contract in the (a,c) plane, a=5.635 Å, b=8.114 Å and c=5.614 Å (c/b=0.69) for xAl Sn =0.25. The decrease of the c/b ratio is in line with the increase of the Goldschmidt tolerance factor t (t = 0.96 for SSO and 0.98 for xAl Sn =0.25) and the progressive orthorhombic to cubic phase transition predicted for large xAl Sn contents. Overall, this study shows that the increase of the valence band width with xAl Sn , which is a desired outcome for p-type TCO materials, results from the substantial increase in Sn-O covalency associated with a negative Moss-Burstein shift (creation of O hole) in the VBM. In summary, new p-type SrSn 1-x Al x O 3 (x = 0, 0.2 and 0.5) perovskites have been successfully synthesized using solid state reaction. The comparison of UV-visible-NIR absorption spectra of SSO and SSAO reveals two important features: 1) an increase in the optical band gap (>3.9 eV) together with the persistence of transparency and 2) the onset of NIR electronic transitions (1400-2700 nm) that increases with xAl Sn . DFT calculations predict a thermodynamic stability domain for p-type charge carriers in SSAOs, which increases with the Sn/Al substitution ratio. The Sn/Al substitution leads to an increase in the Sn-O bond covalency in conjunction with the creation of stable O-holes in the top portion of the valence band. Despite our unsuccessful attempts to carry out electrical conductivity measurements, we believe the structure/property correlation established in this work gives a body of evidence that SSAO materials do behave as p-type transparent conductive oxides. Work is in progress to conclude on that point with thin-film formulation being sought as an alternative to powder synthesis to assess transport properties of SSAO materials. To conclude, chemical substitutions rather than doping give hopes that environmental friendly p-type materials could compete with current n-type materials for TCO applications. The richness of the metal oxide family of compounds provides solid state chemists for an amazing playground to search for the winning combination of elements.

𝑥 2 )

 2 up to 0.22 holes per oxygen.Optical band gaps (𝐸 𝑔 𝑜𝑝𝑡 ) were extracted from the absorption spectra using Tauc's theory39 (see section Methods in ESI) and are displayed in Fig.2b. The extrapolation of the straight line portion of the (αhν)2 plots to α= 0 gives a good estimates of E g opt values. They correspond to 3.9, 4.0 and 4.1 eV for xAl Sn =0, 0.2 and 0.5, respectively, in good agreement with the literature for SSO.40,41

Figure 1 X

 1 Figure 1 X-ray diffraction patterns and enlargement at 2  31° and of SrSn 1-x Al x O 3 (x=0, 0.20 and 0.50) Figure 2 (a) UV-Vis-NIR absorption spectra and (b) optical band gap estimation using Tauc's formula (αh) 2 =f (h) of SrSn0 1-

Figure 3 :

 3 Figure 3: Stability of the oxygen network towards oxygen vacancies V O (V O =  = 0.125) with and without the presence of Sr vacancies V sr (V sr = y = 0.0625). The xAl Sn composition scale is given for the black curve. The Al ratio for the red curve is equal to xAl Sn -2y. The charge carrier type is defined by the sign of the relation (2y+x)-2. Dashed line delimits the n/p-type charge carriers.

Figure 4 :

 4 Figure 4: Band structures computed for (a) SrSnO 3 and (b) SrSn 0.75 Al 0.25 O 3 and their magnification (insets) in the range of 1 eV around the Fermi level (red horizontal line). Fukui functions of (c) SrSnO 3 and (d) SrSn 0.75 Al 0.25 O 3 projected in the ab plane where Sn, Al and O are represented in blue, grey and red, respectively. (e) Fukui function of SrSn 0.75 Al 0.25 O 3 in the mixed Sn/Al and pure Sn planes (ac plane). ELF functions and the corresponding Lewis representations of Sn-O-M (M=Sn or Al) bridge of (g) SrSnO 3 and (h) SrSn 0.75 Al 0.25 O 3 .

Acknowledgements

We thank JL. Bantignies, D. Maurin, A. Grimaud and A. Van De Lee for fruitful technical discussions. A. Berche and P. Jund are also acknowledged for their valuable inputs on physical properties. This work was supported by PHC-Utique (N°3496ZM). Special dedication to Dr Jean-Francois Halet on the occasion of his 60th birthday.

Conflicts of interest

There are no conflicts to declare.

Notes and references