Ruiwen Ji

Meng Lei

Cécile Genevois

Wenda Zhang

Xing Ming

Lunhua He

Mathieu Allix
email: mathieu.allix@cnrs-orleans.fr

Congling Yin
email: congling.yin@glut.edu.cn

Xiaojun Kuang
email: kuangxj@glut.edu.cn

Xianran Xing

Multiple Anion Chemistry for Ionic Layer Thickness Tailoring in Bi 2+2 n O 2+2 n Se n X 2 (X = Cl, Br) van der Waals Semiconductors with Low Thermal Conductivities

published or not. The documents may come

Ⅰ. Introduction

Oxides have attracted extensive interest in solid-state material chemistry, where cationic substitutions are traditionally used to tune the crystal structures and associated physical properties. The anionic substitution, an alternative route to modify the properties and access new mixed-anion compounds, has been largely ignored due to the often-poor stability at ambient atmosphere and the tough synthesis requirements of mixed-anion compounds. [START_REF] Kageyama | Expanding Frontiers in Materials Chemistry and Physics with Multiple Anions[END_REF] Among the known inorganic compounds, numerous inorganic oxides (up to 55781 in the Inorganic Crystal Structure Database, (ICSD, https://icsd.fiz-karlsruhe.de), as of 2 nd May 2021) have been reported, whereas only a limited number of mixed-anion compounds (~3365, i.e. 0.6%) are known, particularly in systems containing ≥ 2 distinct anion species. The crystal structures of oxides are usually described with (pseudo) close-packed arrangements of O 2-anions (and similar sized cations with O 2-) and smaller-sized cations filling the resulting voids, for which the typical structures include A3O4 spinel, AO rock-salt, and ABO3 perovskite based on the common coordination polyhedra as building units (e.g. AO4 tetrahedron, AO6 octahedron, and AO12 cuboctahedron). However, for anionic substitution, as most of the anions are larger (and more polarizable, or softer) than O 2-e.g. Cl -, Se 2-and As 3-, they tend to aggregate two-dimensionally to minimally mix with smaller sized (and harder) O 2- layers to save space, and destroy the 3-dimensional (3D) close packing of the anions. Distinct 2dimensional (2D) structures are thus generated and exhibit rich physical properties including the technologically-important high TC superconductivity (e.g. LaOFeAs), 2 thermoelectricity (e.g. BiOCuSe and BiOCuTe), [START_REF] Zhou | Promising Thermoelectric Bulk Materials with 2D Structures[END_REF][START_REF] Zhao | BiCuSeO Oxyselenides: New Promising Thermoelectric Materials[END_REF][START_REF] Vaqueiro | A Copper-Containing Oxytelluride as a Promising Thermoelectric Material for Waste Heat Recovery[END_REF] transparent p-type semiconductivity (e.g. LaOCuS), [START_REF] Ueda | Transparent p-Type Semiconductor: LaCuOS Layered Oxysulfide[END_REF] ionic conductivity (i.e. LaOAgS and LaOCl), [START_REF] Imanaka | Water-Insoluble Lanthanum Oxychloride-Based Solid Electrolytes with Ultra-high Chloride Ion Conductivity[END_REF][START_REF] Wilmer | Two-Dimensional Silver Fast-Ion Conduction in (LaO)AgS[END_REF] and photocatalysis (BiOCl and Bi4NbO8Cl). [START_REF] Fujito | Layered Perovskite Oxychloride Bi4NbO8Cl: A Stable Visible Light Responsive Photocatalyst for Water Splitting[END_REF][START_REF] Jiang | Synthesis and Facet-Dependent Photoreactivity of BiOCl Single-Crystalline Nanosheets[END_REF] Multiple anion compounds have been standing at the frontier of solid-state materials chemistry and physics. [START_REF] Kageyama | Expanding Frontiers in Materials Chemistry and Physics with Multiple Anions[END_REF][START_REF] Harada | Heteroanionic Materials by Design: Progress toward Targeted Properties[END_REF] Up to now, the exploration of multiple-anionic systems is mainly limited to several binary anionic systems, such as oxyhydrides, oxyhalides, and oxychalcogenides. Among these binary-anionic systems, layered structures containing the common structural motif of fluorite-type A2O2 2+ (A = Ln 3+ , Bi 3+ ) layers with a significant contribution of covalent interactions are attracting more and more attention, e.g. Bi2O2Se, [START_REF] Wu | High Electron Mobility and Quantum Oscillations in Non-Encapsulated Ultrathin Semiconducting Bi2O2Se[END_REF] Bi2O2Cl2, 10 Bi2O2Cu2Se2, [START_REF] Mcclain | In Situ Mechanistic Studies of Two Divergent Synthesis Routes Forming the Heteroanionic BiOCuSe[END_REF] and La2O2Fe2OSe2. [START_REF] Free | Low-Temperature Nuclear and Magnetic Structures of La2O2Fe2OSe2 from X-Ray and Neutron Diffraction Measurements[END_REF] In these layered multiple anionic compounds, the fluorite-type A2O2 2+ layers, which are constructed from edge-sharing OA4 tetrahedra, alternate with the charge-balancing anionic layers. In the simplest case, the charge-balancing anionic layers are single-layered Ch 2-anions (Figure 1a), as shown in Bi2O2Ch (Ch 2-= S 2-, Se 2-, Te 2-, or NCN 2-), 12, 15-17 and double-layered X -anions (Figure 1b) in the case of Bi2O2X2 (or BiOX equally, X = Cl -, Br -, or I -). [START_REF] Jiang | Synthesis and Facet-Dependent Photoreactivity of BiOCl Single-Crystalline Nanosheets[END_REF][START_REF] Lotnyk | Influence of the Reduced Dimensionality on the Thermodynamical and Electrical Properties of Photosensitive BiOX (X = Cl, Br, and I) Semiconductors[END_REF][START_REF] Bhachu | Bismuth Oxyhalides: Synthesis, Structure and Photoelectrochemical Activity[END_REF] Filling of the tetrahedral voids between the double anionic layers in these structures generates antifluorite-type M2Pn2 2-layers forming quaternary A2O2M2Pn2 materials (A = La 3+ or Bi 3+ , M 2+ = Fe 2+ or Mn 2+ , Pn 3-= As 3-, P 3-, or Sb 3-) with ZrCuSiAs-like structures (Figure 1c). [START_REF] Kamihara | Iron-Based Layered Superconductor La[O1-xFx[END_REF][START_REF] Mcclain | In Situ Mechanistic Studies of Two Divergent Synthesis Routes Forming the Heteroanionic BiOCuSe[END_REF][START_REF] Nientiedt | Quaternary Equiatomic Manganese Pnictide Oxides AMnPO (A = La-Nd, Sm, Gd-Dy), AMnAsO (A = Y, La-Nd, Sm, Gd-Dy, U), and AMnSbO (A = La-Nd, Sm, Gd) with ZrCuSiAs Type Structure[END_REF] In the case of A2O2M2OSe2 materials (A= La 3+ -Sm 3+ , M 2+ = Fe 2+ or Mn 2+ ), Se 2-anionic double layers embedded by an M2O 2+ planar layer of anti-CuO2 type in between are incorporated between the A2O2 2+ cationic layers to compensate the positive charge, [START_REF] Free | Low-Temperature Nuclear and Magnetic Structures of La2O2Fe2OSe2 from X-Ray and Neutron Diffraction Measurements[END_REF][START_REF] Free | Synthesis, Structure and Properties of Several New Oxychalcogenide Materials with the General Formula A2O2M2OSe2 (A = La-Sm, M = Fe, Mn)[END_REF] as shown in Figure 1d. In these above-mentioned layered structures, Bi2O2Se, Bi2O2Cu2Se2, and La2O2Fe2OSe2 are ionic layered materials, featuring alternative stacking of cationic and anionic layers, forming two-dimensional covalent-dominant structural units which are connected by ionic bonding interactions in the third dimension. Interestingly, the Bi2O2X2 (X = Cl -, Br -, or I -) structure is composed of cationic fluorite Bi2O2 2+ layers bound on either side to terminal X -anion layers in the sequence X-Bi2O2-X-X-Bi2O2-X, forming a van der Waals gap in the successive halide layers. Thus, this material can be described as layered XBi2O2X slabs stacked by X•••X van der Waals interactions in the third dimension, and allow easy exfoliation or deposition as atomically thin nanosheets, [START_REF] Zhao | Oxygen Vacancy Dependent Photocatalytic CO2 Reduction Activity in Liquid-Exfoliated Atomically Thin BiOCl Nanosheets[END_REF][START_REF] Nalawade | All-Printed Dielectric Capacitors from High-Permittivity, Liquid-Exfoliated BiOCl Nanosheets[END_REF] which is useful in the fabrication of functional heterostructures and devices, spanning electronic to energy materials. Thus, among these chargeredeeming layers inserted between A2O2 2+ layers, the double-layer X -anions are most spectacular in structural modification and can act as scissors cutting the 3D ionic layered structures into 2D van der Waal slabs, leading to new van der Waals materials with highly anisotropic physical properties, as proposed in the recently reported multiple anion material La5O4Cu4As4Cl2 along with a P-containing analogue. [START_REF] Bartsch | La3Cu4P4O2 and La5Cu4P4O4Cl2: Synthesis, Structure and 31 P Solid State NMR Spectroscopy[END_REF][START_REF] Eul | An Extension of Pnictide Oxide Chemistry -Salt Flux Synthesis and Structure of La5Cu4As4O4Cl2[END_REF] Given the common fluorite-type A2O2 2+ structural building units in these binary anionic layered structures, it would be expected that more complex structures may be generated through intergrowth of these structure types, thus allowing tailoring the physical properties when three or more anions are included in a single structure. However, these compounds are rarely explored, and there is very limited research on such multiple-anionic compounds. Some of them are isoelectronic and isostructural to materials with two-anion compounds, e.g. (SrF)2Fe2OSe2 26 is isostructural with (LaO)2Fe2OSe2. [START_REF] Free | Low-Temperature Nuclear and Magnetic Structures of La2O2Fe2OSe2 from X-Ray and Neutron Diffraction Measurements[END_REF] The direct combination of Bi2O2Se and Bi2O2Cu2Se2 slabs remained unsuccessful but was recently circumvented by the addition of a third halide anion, leading to a series of multiple anionic Bi2+2nO2+2nCu2-δSe2+n-δXδ (X = Cl and Br, n =1 and 2) defective-layer materials. [START_REF] Gibson | Bi2+2nO2+2nCu2-δSe2+n-δXδ (X = CI, Br): A Three-Anion Homologous Series[END_REF][START_REF] Gibson | Bi4O4Cu1.7Se2.7Cl0.3: Intergrowth of BiOCuSe and Bi2O2Se Stabilized by the Addition of a Third Anion[END_REF] The partial substitution of the X -anions for Se 2-in the Bi2O2Se slab and the resulting charge-compensating Cu vacancies in the Bi2O2Cu2Se2 layers are of vital importance in the stabilization of the Bi2+2nO2+2nCu2-δSe2+n-δXδ series. These intergrowth materials showed exceptionally low thermal conductivity (0.4 W/mK and lower) and narrow bandgaps (< 0.5 eV), compared with both parent compounds, [START_REF] Zhao | BiCuSeO Oxyselenides: New Promising Thermoelectric Materials[END_REF][START_REF] Tan | Enhanced Thermoelectric Performance of n-Type Bi2O2Se by Cl-Doping at Se Site[END_REF][START_REF] Zhan | High-Temperature Thermoelectric Behaviors of Sn-Doped n-Type Bi2O2Se Ceramics[END_REF][START_REF] Ruleova | Thermoelectric Properties of Bi2O2Se[END_REF] making them potential thermoelectric materials. [START_REF] Gibson | Bi2+2nO2+2nCu2-δSe2+n-δXδ (X = CI, Br): A Three-Anion Homologous Series[END_REF][START_REF] Gibson | Bi4O4Cu1.7Se2.7Cl0.3: Intergrowth of BiOCuSe and Bi2O2Se Stabilized by the Addition of a Third Anion[END_REF] Later, another novel layered structure Bi4O4SeCl2 was stabilized as a 1:1 superlattice of the structural units present in the van der Waals insulator Bi2O2Cl2 and the ionic layered semiconductor Bi2O2Se. This compound can be noted as (m = 1, n = 1) alternative stacking between Bi2O2Cl2 and Bi2O2Se units (Figure 2a), respectively, featuring partial anionic disorder of Se and Cl between the Bi2O2 2+ cationic layers. The van der Waals semiconductor Bi4O4SeCl2 retains the high and anisotropic electronic mobility of Bi2O2Se [START_REF] Wu | High Electron Mobility and Quantum Oscillations in Non-Encapsulated Ultrathin Semiconducting Bi2O2Se[END_REF][START_REF] Chen | Electronic Structures and Unusually Robust Bandgap in an Ultrahigh-Mobility Layered Oxide Semiconductor[END_REF] while reducing the dimensionality of the ionic bonding network of Bi2O2Se to allow facile exfoliation of Bi4O4SeCl2 to 1.4 nm height. [START_REF] Gibson | Modular Design via Multiple Anion Chemistry of the High Mobility van der Waals Semiconductor Bi4O4SeCl2[END_REF] Due to the anisotropic bond contrast in vdW Bi2O2Cl2 slab and the interface mismatch (between the Bi2O2 2+ and Se 2-layers) in the Bi2O2Se slab, the contribution of the longitudinal and transverse phonons to heat transport are significantly suppressed in the layered materials Bi4O4SeCl2, leading to an extremely low thermal conductivity of 0.1W/mK at room temperature along the c-axis direction, which is about only four times the thermal conductivity of air. 34 In Bi4O4SeCl2, the stacking arrangement of Bi2O2Se and Bi2O2X2-type units is not well understood and questions remain regarding the possible control of the thicknesses of the ionic-bonding layers for tuning the exfoliated nanosheet heights for structure fabrication and associated physical properties. Here, based on the dimensionality tailoring effect of double-layer halogen anions in the Bi2O2Se network into van der Waals slabs, we expand the Bi2O2X2-Bi2O2Se intergrowth of van der Waals compounds from the reported compound Bi4O4SeCl2 to the homologous series Bi2+2nO2+2nSenX2 (X = Cl and Br) via stoichiometry tuning, building a series of van der Waals materials with different thicknesses of ionic-bonding layers, as illustrated in Figures 2a-c, which demonstrates new insight into multiple anion chemistry for tailoring the ionic-bonding network into different thicknesses in the layered structures.

Ⅱ. Experimental Procedure

Synthesis. The polycrystalline samples of Bi6O6SeCl4, Bi4O4SeX2 (X = Cl and Br), Bi6O6Se2X2 (X = Cl and Br), and Bi8O8Se3Cl2 were prepared by solid-state reaction method. The starting materials of Bi (99.99%, Aladdin), Bi2O3 (99.9%, Aladdin), Se (99.9%, Aladdin), and lab-made reagents BiOX (X = Cl and Br, see details below) were weighed according to the correct stoichiometric ratios and thoroughly mixed in an agate mortar with a pestle. The mixtures were pressed into pellets of ~ 6 mm in diameter and ~ 2 mm in thickness under a uniaxial pressure of ~150 MPa and transferred to alumina crucibles, and then sealed into an evacuated quartz tube, which was fired at 700 °C for 48 hours with one intermediate pressing and grinding, and finally slowly cooled down to room temperature at 5 °C/min. The experimental densities of the ceramic pellets were calculated using the measured geometric sizes (diameters and thicknesses) and the masses of the pellets, while the theoretical densities were calculated based on the chemical compositions and refined lattice parameters by Rietveld methods. The obtained dark pellets of Bi4O4SeX2 and Bi6O6Se2X2 (X = Cl and Br) reached ~ 80% of the X-ray theoretical densities.

The BiOX (X = Cl and Br) precursors were synthesized using an hydrothermal reaction 10 from the starting materials of Bi(NO3)3•5H2O (99.9%, Aladdin), NaCl (99.5%, Xilong Scientific), and KBr (99.5%, Aladdin) and their phase purities were confirmed by powder X-ray diffraction (XRD). All the materials used and obtained here are air-stable and safe to be handled in ambient conditions.

Characterization. Phase formations were examined by powder XRD at ambient conditions, which was conducted on a PANalytical X'Pert Pro diffractometer in Bragg-Brentano geometry using Cu Kα radiation.

The samples were dispersed on a zero-background Si substrate using a drop of ethanol to reduce preferred orientation because of its layered structural feature. High-quality laboratory XRD data for Rietveld refinements were collected over the 2θ range of 10-120°, using sample spinning mode. Time-of-flight (TOF) neutron powder diffraction (NPD) data were collected on the General-Purpose Powder Diffractometer (GPPD) at China Spallation Neutron Source (CSNS). Ab initio structure solutions for Bi6O6Se2X2 were carried out in EXPO2014, using the direct space method. [START_REF] Altomare | A Kit of Tools for Phasing Crystal Structures from Powder Data[END_REF][START_REF] Altomare | Structure Solution by Powder Data in Direct and Reciprocal Space[END_REF] Rietveld refinements were performed against XRD and NPD data utilizing the Topas Academic 6 program. [START_REF] Coelho | Topas Academic Version 5[END_REF] Bond valence sums (BVSs) were calculated by Brown and Altermatt's method. [START_REF] Brown | Bond-Valence Parameters Obtained from a Systematic Analysis of the Inorganic Crystal-Structure Database[END_REF] Selected-area electron diffraction (SAED) and convergent-beam electron diffraction (CBED) patterns were recorded on a JEOL JEM-2100F transmission electron microscope (TEM) with an accelerating voltage of 200 kV. High-resolution scanning transmission electron microscopy (STEM) -high-angle annular dark-field (HAADF) / annular bright-field (ABF) imaging was performed on a JEOL ACCELARM 200 cold FEG TEM operating at 200 kV, equipped with double spherical aberration (Cs) correctors. The STEM-HAADF/ABF images were acquired with an inner-outer collection angle of 90-370mrad / 16-32mrad respectively and a 0.1 nm probe size. The sample was first crushed and dispersed in ethanol, and one drop of the solution with the small crystallites in suspension was deposited onto a carbon-coated copper grid. Before the observations, the grid was cleaned for 1 minute with a plasma to prevent carbon contamination. Scanning electron microscopy (SEM) was performed with a GeminiSEM 300 (ZEISS, Germany) on the polished surface of the ceramic sample, which was deposited on a carbon tape attached to an aluminum stub and coated with a thin film of gold before performing imaging.

The optical diffuse reflectance data were collected in the 280-2500 nm range using a Lambda 750 PerkinElmer spectrometer equipped with an integrating sphere. AC impedance spectroscopy measurements were performed on a Solartron 1260 frequency response analyzer over a frequency range of 10 -1 -10 7 Hz within 7-300 K. The silver paste was coated on the opposite faces of the pellets, which were then clamped with the cupper pads in the Physike SHI-4-2 cryostat. An oil-free Edwards T-station 85 molecular pump was used to produce a vacuum below 10 -4 mbar in the cryostat. Then, a Sumitomo HC-4E compressor was employed to cool the system to the desired temperatures through a Physike C22 temperature controller. The impedance data analysis was carried out through the equivalent circuit fitting using ZView software. [START_REF] Johnson | ZView for Windows, Impedance/Gain phase graphing and Analysis Software[END_REF] High precision thermal conductivity of pellets was measured perpendicular to the pellet plane in the 70-300 K temperature range by a two-probe method under high vacuum conditions on the homemade Cryogenic Property Measurement System (CPMS-77). The thermal conductivity data have ~ 4% standard deviation based on the calibration against the stainless-steel standard sample. [START_REF] Zhou | Design of a Device for Measuring Thermal Conductivity of Solid Materials with Miniature Refrigerator as Cold Source[END_REF] Theoretical calculations. Geometry optimization and electronic structure calculations were performed by using the Vienna Ab initio Simulation Package (VASP) based on density functional theory (DFT). [START_REF] Kresse | Efficient Iterative Schemes for Ab Initio Total-Energy Calculations Using a Plane-Wave Basis Set[END_REF][START_REF] Kresse | Efficiency of Ab-Initio Total Energy Calculations for Metals and Semiconductors using a Plane-Wave Basis Set[END_REF] The projector-augmented-wave potential and the Perdew-Burke-Ernzerhof (PBE) exchangecorrelation functional were employed. [START_REF] Perdew | Generalized Gradient Approximation Made Simple[END_REF] The cutoff energy of 520 eV was specified for the plane-wave basis set, and the smallest allowed spacing of the reciprocal space between k points was set by the KSPACING parameter with a value of 0.2 Å -1 . The conjugate gradient algorithm was used to perform the structural energy minimization, and the convergence criteria were set to 10 -6 eV and 0.01 eV/Å for total energy and atomic force, respectively. The DFT-D2 method of Grimme was adopted to describe the van der Waals interactions of the layered materials. [START_REF] Grimme | Semiempirical GGA-Type Density Functional Constructed with a Long-Range Dispersion Correction[END_REF] According to the experimental data of the studied materials, the ideal structure models without atomic mixed-occupancy were built and further optimizations were carried out to relax the atomic positions and lattice constants. As shown in Table S1, the theoretically optimized lattice constants are consistent with our experimental determined values, demonstrating the validity of the calculation parameters.

Ⅲ. Results

Crystal Structures of Bi4O4SeX2 (X = Br and Cl). The crystal structure of Bi4O4SeCl2 was previously reported based on single-crystal X-ray diffraction by Gibson et al. [START_REF] Gibson | Modular Design via Multiple Anion Chemistry of the High Mobility van der Waals Semiconductor Bi4O4SeCl2[END_REF] However, because the X-ray scattering length is proportional to the atomic number (Z) of the elements, the XRD intensity of the Bi4O4SeCl2 data is dominated by the bismuth atoms contribution. The lighter elements Se (Z =34), Cl (Z = 17) and O (Z = 8) each contribute to only ~7.5% of the total electron density per formula and thus XRD intensity, resulting in low-intensity statistics insufficient for accurate determination of the occupancies and positions on these anionic. The information for these light elements can be improved using neutron diffraction since the neutron coherent scattering lengths for Cl, Se and Bi are of same order (9.577, 7.970, and 8.532 fm respectively) and even the lightest O has a decent neutron scattering length (5.083 fm). [START_REF] Sears | Neutron Scattering Lengths and Cross Sections[END_REF] .

In this study, Rietveld refinements of Bi4O4SeCl2 were carried out against both NPD and XRD data simultaneously (as shown in Figures 3a and3b), starting from the previously reported X-ray structure.

To retain the nominal stoichiometry and charge neutrality during the refinement, the Cl occupancy on the two-fold Se1 site in the Bi2O2Se layers was set as twice the amount of Se substitution rate on the fourfold Cl2 site in the Bi2O2Cl2 layers, subject to the overall occupancies of either anionic site being unity.

Good fits were obtained with an overall weighted profile Rwp factor of 5.78 % and goodness of fit χ 2 of 1.17. Synthesis of the Br-substituted analogue was also attempted at similar synthetic conditions and a pure phase sample of composition Bi4O4SeBr2 was identified with a similar XRD pattern (Figure 4a) as that of Bi4O4SeCl2. The crystal structure of Bi4O4SeBr2 was refined simultaneously against the XRD and NPD data (Figures 4a and4b), using the Bi4O4SeCl2 structure as the starting model (Figure 2a). Here neutron diffraction is essential to discern the Se and Br atoms Bi4O4SeBr2 as they are indistinguishable in the Xray data. The coherent scattering lengths for Se and Br are 7.97 and 6.795 fm, respectively, leading to a small but existing contrast between these two atoms in the NPD data refinement. The site occupancies of Se and Br atoms were refined under the similar constraints used for the Cl-compound case, which leads to a Br substitution rate on the Se1 site ~ 0.41(4) and a Se substitution rate on the Br2 site ~ 0.21 (2). Good Rietveld refinements are obtained with an overall Rwp of 6.32% and χ 2 of 1.74. The final refined structural parameters and selected interatomic distances in Bi4O4SeBr2 are given in Table 2. Bi4O4SeBr2 is isostructural with Bi4O4SeCl2 (Figure 2a), built with alternating Bi2O2Se and Bi2O2Br2 type layers. [START_REF] Shannon | Revised Effective Ionic Radii and Systematic Studies of Interatomic Distances in Halides and Chalcogenides[END_REF] and the layered structural anisotropy. However, the difference between the c-lattice parameter of Bi4O4SeX2 (X = Br and Cl), as deduced from the sum of one c-lattice parameter of Bi2O2Se and two c-lattice parameters of Bi2O2X2, should be even larger ~1.50 Å, supposing no Se/X disorder occurred in Bi4O4SeX2 and its Bi2O2Se and Bi2O2X2 slabs were identical with those in the end members. Thus, the reduced c-periodicity difference between the Cl and Br analogue keeps with Se/X anionic disorder in the relaxed Bi2O2Se and Bi2O2X2 slab, which leads to a more compacted layer stacking and helps to stabilize the layered structures of Bi4O4SeX2. Furthermore, compared with the ionic layered compounds Bi4O4Cu1.7Se2.7X0.3 having 30% Se atoms replaced by halide in the Bi2O2Se layer, [START_REF] Gibson | Bi2+2nO2+2nCu2-δSe2+n-δXδ (X = CI, Br): A Three-Anion Homologous Series[END_REF][START_REF] Gibson | Bi4O4Cu1.7Se2.7Cl0.3: Intergrowth of BiOCuSe and Bi2O2Se Stabilized by the Addition of a Third Anion[END_REF] the X substitution rates on the Se site in the Bi2O2Se layer reach ~50% in the van der Waals materials Bi4O4SeX2, i.e. 0.55(1) and 0.41(4) for the X = Cl -and Br -compounds. This results in a high content of substituted Se defects by X anions, close to the Se vacancies level ~50% in the well-relaxed surface layers of Bi2O2Se crystal. [START_REF] Chen | Electronic Structures and Unusually Robust Bandgap in an Ultrahigh-Mobility Layered Oxide Semiconductor[END_REF] This observation indicates that the weak interlayer vdW interactions in the Bi2O2Cl2 slab allow nearly free relaxation of Bi2O2Se layers in the ab plane in the vdW materials Bi4O4SeX2, while the strong bond interaction in the Bi2O2Cu2Se2 slab limits the relaxation of Bi2O2Se layers in the Bi4O4Cu1.7Se2.7X0.3 materials. The ab initio structure solution for Bi6O6Se2Cl2 based on the XRD data was then performed in the space group P4/nmm by the direct method using EXPO 2014, [START_REF] Altomare | A Kit of Tools for Phasing Crystal Structures from Powder Data[END_REF][START_REF] Altomare | Structure Solution by Powder Data in Direct and Reciprocal Space[END_REF] in which the initial positions of heavier bismuth and selenium atoms can be directly determined. Subsequently, the positions of the lighter oxygen and chlorine atoms were obtained by difference Fourier analysis. Rietveld refinements of lattice parameters, atomic positions, and isotropic displacement parameters (Biso) were carried out simultaneously against XRD and NPD data. Similarly to the Bi4O4SeX2 case, neither oxygen vacancy nor anionic disorder was observed on the O1 and O2 sites in the Bi2O2 2+ layers, while the anionic mixing of Se and Cl on the Se1 and Cl2 sites were identified. Therefore, the Cl occupancy on the Se1 site was set identical with the Se occupancy on the Cl2 site as the Se1 and Cl2 sites have the same multiplicity (Table 3), and the overall occupancies of both Se1 and Cl2 sites are subject to unity in the following refinements.

Synthesis of

The final refinement converged to Rwp of 7.16% and χ 2 of 2.6, and the Rietveld plots of XRD and NPD data are shown in Figure 6. The final refined atomic coordinates and selected bond distances are given in Table 3.

The structure of Bi6O6Se2Cl2 can be considered as a 2:1 intergrowth between Bi2O2Se and Bi2O2Cl2 slabs (corresponding to the same composition parent compounds ), resulting in a -Bi2O2-Se-Bi2O2-Se-Bi2O2-Cl-Cl-layer sequence (Figure 2b), in contrast to the (m = 1, n = 1) phase Bi4O4SeCl2 showing a stacking sequence of -Bi2O2-Se-Bi2O2-Cl-Cl-(Figure 2a). The successive Cl -layers introduce the van der Waals gap and prevent further extension of ionic bonding interactions along the c axis in Bi6O6Se2Cl2.

Therefore, the Bi6O6Se2Cl2 structure can be represented as ideally charge-neutral Bi6O6Se2Cl2 slabs of 1.96 nm thickness, stacked by van der Waals forces along the c direction.

The Bi6O6Se2Cl2 slab (Figure 7a) can be seen as a Bi4O4SeCl2 slab intergrown with a Bi2O2Se slabs and the van der Waals interaction that stacks these slabs into three-dimensional crystals. Anionic disorder significantly decreases the internal energy of the Bi4O4SeCl2 (and Bi6O6SeCl2) vdW slabs in the following three aspects.

First, the Slater atomic radius ratio f = (RA)/(RX + 0.2RM) of Bi2O2Se, which was defined for the AM2X2 structure (here A = Se, M = Bi, and X = O), is 1.25, which places Bi2O2Se on the phase boundary between the trigonal 6-coordinated anti-CaAl2Si2 structure (Figure S3a) and tetragonal 8-coordinated anti-ThCr2Si2 structure (Figure S3b). [START_REF] Peng | Crystal Chemistry and Thermoelectric Transport of Layered AM2X2 Compounds[END_REF] This suggests that Se is undersized in its eight-coordinate environment (with a long 3.4 Å Bi-Se bond) through mismatch to the Bi2O2 layer dimensions. The internal lattice strain from this bonding frustration at the bridging Bi2O2-Se interface positions the tetragonal structure adjacent to a distorted phase [START_REF] Wu | Bismuth Oxychalcogenides: A New Class of Ferroelectric/Ferroelastic Materials with Ultra High Mobility[END_REF] that is achieved through in-plane static displacements in orthorhombic ultrathin, freestanding Bi2O2Se nanosheets. [START_REF] Ghosh | Ultrathin Free-Standing Nanosheets of Bi2O2Se: Room Temperature Ferroelectricity in Self-Assembled Charged Layered Heterostructure[END_REF] This phase stability of the tetragonal Bi2O2Se slab versus the distorted orthorhombic Bi2O2Se slab can be enhanced by the introduction of Se deficiency via chemical substitution of Cl -on the Se site.

Second, the parent compounds Bi2O2Se and Bi2O2Cl2 have slightly different Bi2O2 2+ layers due to the coordination requirement of Se 2-and Cl -in the charge-balancing layers, as reflected by their tetragonal lattice parameter a = 3.897(2) Å and 3.8870(5) Å respectively. [START_REF] Mao | Synthesis and Electrical Transport Properties of Bi2O2Se Single Crystals[END_REF][START_REF] Keramidas | The Crystal Structure of BiOCl[END_REF] Thus, in the hypothetical perfectly ordered Bi4O4SeCl2 (and Bi6O6Se2Cl2) structure, the sandwiched Bi2O2 2+ layers between the Bi2O2Se and Bi2O2X2 intergrowth layers are thus strained due to the different requirement of Se and Cl ions on either side. This internal lattice strain from the different bonding requirements of Se and Cl ions at the bridging Bi2O2-Se interface and terminal Bi2O2-Cl interface can be partially released by the mutual anionic substitution of Se and Cl on both sides.

Third, the anion disorder will break the charge neutrality of bridging Bi2O2Se layers, and introduce the charge fluctuation in the terminal Bi2O2X2 layers. Opposite charges for these two intergrowth layers and thus electrostatic interactions between the two slabs to further decrease the total energy.

Based on the observations on the anionic site disorder and its importance on the stacking stability of As STEM-ABF image intensity I is proportional to Z 1/3 , the ABF image contrast is more sensitive to light elements than the STEM-HAADF image. [START_REF] Findlay | Robust Atomic Resolution Imaging of Light Elements Using Scanning Transmission Electron Microscopy[END_REF] Thus, STEM-ABF images were collected on the same zone to localise oxygen layers. In contrast with bright spheres on the STEM-HAADF image, the atoms on the STEM-ABF image appear as black dots. The large black dots correspond to the Bi arrays, and the intermediate ones correspond to Se and Cl arrays, as shown in Figure 8g right. The small black dots pointed by red arrows correspond to the oxygen layers. It should be noted that even on the ABF image the signal from oxygen atoms is still very weak due to its localization at short distances from much heavier Bi cations. By coupling both HAADF and ABF imaging modes, the entire atomic position information in the structure can be accessible (Figure 8g the intrinsic absorption edge of 860-1000 nm. [START_REF] Sun | Bismuth-Based Halide Double Perovskite Cs2LiBiCl6: Crystal Structure, Luminescence, and Stability[END_REF][START_REF] Fox | Optical Properties of Solids Oxford University Express[END_REF] From the corresponding Tauc plots (Figure 9b), the indirect bandgaps of Bi4O4SeBr2, Bi4O4SeCl2, and Bi6O6Se2Cl2 were estimated to be 1.07(5), 1.05(5), and 1.04 (5) eV respectively, the value for Bi4O4SeCl2 matching well with the previously reported value of 1.15(5) eV. [START_REF] Gibson | Modular Design via Multiple Anion Chemistry of the High Mobility van der Waals Semiconductor Bi4O4SeCl2[END_REF] An extrinsic defect-related sub-bandgap was estimated to be 0.73 (5) and 0.75(5) eV for Bi4O4SeBr2 and Bi4O4SeCl2 respectively (Figure 9b). crystals along the c axis can be estimated to be about 10 2 Ω cm, while the in-plane resistivity can be estimated to be ~ 10 -2 Ω cm as referenced from previous Bi4O4SeCl2 single crystal measurements. [START_REF] Gibson | Modular Design via Multiple Anion Chemistry of the High Mobility van der Waals Semiconductor Bi4O4SeCl2[END_REF] The high anisotropic resistivity showing about 4 orders of magnitude difference along and perpendicular to the ab plane is in line with the layered crystal structure nature. Similar anisotropic conductivity is also observed in other 2D vdW material systems, such as graphite. [START_REF] Greenwood | Chemistry of the Elements[END_REF] The total resistivity of these four samples (Figure 12b) shows clear decreasing trends with temperature, indicating the overall nondegenerate semiconductors which carrier concentration dominates the resistivity over mobility in the grain boundary. (b) (a) pattern (dashed red line), showing a strong increase in the intensity of (00l) peaks and a reduction of (hk0) peaks (highlighted in the inset), indicating a strong preferred orientation of the [001] direction.

Thermal conductivity measurements performed parallel to the pressing axis of the pellet on all compounds showed very low thermal conductivities at room temperature, with values of 0.66(3), 0.58(2), 0.33(1), and 0.13(1) W/mK for Bi6O6Se2Cl2, Bi6O6Se2Br2, Bi4O4SeCl2, Bi4O4SeBr2, respectively (Figure 14). The value of our Bi4O4SeCl2 pellet is larger than the reported value of 0.10(2) W/mK measured on a (001) textured Bi4O4SeCl2 pellet obtained using spark plasma sintering (SPS). We stipulate that this difference is related to the fact that our pellet sample is less (001) textured and therefore shows more contribution from the in-plane thermal conductivity of 1.0 W/mK. Nevertheless, all compounds exhibit exceptionally low thermal conductivities for crystalline compounds, as well as an unusual temperature dependence. Because the pellets have similar density and extent of (001) texture (Figures 14 andS7-S9), their thermal conductivity can be compared. The thermal conductivity and total electronic conductivity of the four samples above 100 K have similar trends and decrease from Bi4O4SeBr2, Bi4O4SeCl2, Bi6O6Se2Br2 to Bi6O6Se2Cl2. This observation indicates that the heat transport of Bi2+nO2+nSeXn materials is probably mediated not only by photons scattering but also by electron scattering as observed in metals and alloys given their high carrier concentration. Compared with the Bi4O4SeX2 (X = Br and Cl) materials, the Bi6O6Se2X2 (X = Br and Cl) pellets exhibit higher thermal conductivities, which can be linked to the higher proportion of the Bi2O2Se slab. The Bi2O2Se slab allows much larger thermal conductivity along the c axis than the Bi2O2X2 slab with strong bond contrast, as revealed by the Bi2O2Se and Bi2O2Cl2 pellets demonstrating a thermal conductivity of ~ 1.25 and 0.25 W/mK respectively. [START_REF] Gibson | Modular Design via Multiple Anion Chemistry of the High Mobility van der Waals Semiconductor Bi4O4SeCl2[END_REF] Furthermore, for Bi4O4SeX2 and Bi6O6Se2X2 structures, the Br-containing materials show lower thermal conductivity than their Cl analogues. A similar observation has been reported for the Bi4O4Cu1.7Se2.7X0.3 and Bi6O6Cu1.6Se3.6X0.4 layered materials. [START_REF] Gibson | Bi2+2nO2+2nCu2-δSe2+n-δXδ (X = CI, Br): A Three-Anion Homologous Series[END_REF][START_REF] Gibson | Bi4O4Cu1.7Se2.7Cl0.3: Intergrowth of BiOCuSe and Bi2O2Se Stabilized by the Addition of a Third Anion[END_REF] The phenomenon can be related to the fact that Cl possesses stronger electron affinity and lighter atomic weight than Br. Therefore, compared with the Br analogues, the Cl-containing materials have larger internal energy E, as evidenced by higher Debye temperature of BiOCl than BiOBr, [START_REF] Lotnyk | Influence of the Reduced Dimensionality on the Thermodynamical and Electrical Properties of Photosensitive BiOX (X = Cl, Br, and I) Semiconductors[END_REF] and lower density ρ (Table 1-3), which leads to a larger photon velocity v according to the v = (E/ρ) 0.5 and an increased photon-mediated thermal conductivity.

Conclusion

A series of multiple-anionic vdW materials Bi2+2nO2+2nSenX2 (n = 1-3, X = Cl, Br) were synthesized using solid-state reactions, leading to high-order layers with variable ionic Bi2O2Se layer thicknesses through the Bi2O2Cl2 cutting effect. Bi4O4SeX2 and Bi6O6Se2X2 were isolated as single-phase materials, showing ~50% Se replaced by X anions in the Bi2O2Se slab as a common feature, while the other stackings appear to be mixed phase. The Bi8O8Se3Cl2 composition leads to an interesting mixed phase containing different and complex intergrowth structures in crystal single crystallite. The Se/X anionic disorder in the Bi2O2Se and Bi2O2Cl2 slabs plays an important role in the stabilization of the Bi4O4SeX2 and Bi6O6Se2X2 polymorphs, which reduces the lattice strain on the sandwiched Bi2O2 interface and introduces charge separation on Bi2O2Se and Bi2O2X2 layers, allowing for the combination of different structural motifs. All

Figure 1 .

 1 Figure 1. Typical layered structures with two anions: (a) Bi2O2Se, (b) Bi2O2Cl2, (c) Bi2O2Cu2Se2, and (d)

Figure 2 .

 2 Figure 2. Comparison of the crystal structures of (a) Bi4O4SeCl2, (b) Bi6O6Se2Cl2, and (c) Bi8O8Se3Cl2

Figure 3 .

 3 Figure 3. Rietveld plots of (a) XRD and (b) NPD data for Bi4O4SeCl2. The inset in (a) enlarges the fit to

Figure 4 .

 4 Figure 4. Rietveld plots of (a) XRD and (b) NPD data for Bi4O4SeBr2. The inset in (a) enlarges the fit to

  Higher-Order Homologues. In order to expand the search to higher-order homologues, synthesis attempts were carried out for (m = 1, n = 2), (m = 1, n = 3), and (m = 2, n = 1) stackings, corresponding to the Cl-based compositions Bi6O6Se2Cl2, Bi8O8Se3Cl2, and Bi6O6SeCl4, respectively. No evidence of the (m = 2, n = 1) phase was found in the synthesis attempts of Bi6O6SeCl4, which produce a mixed (m = 1, n = 1) phase Bi4O4SeCl2 and Bi2O2Cl2 as the final product, as evidenced by the Rietveld XRD plot (FigureS1). In contrast, the synthesis endeavors for (m = 1, n = 2) phase at the composition of Bi6O6Se2Cl2 produced a single-phase compound, and its structural analysis is provided in the following sections. The synthesis of (m = 1, n = 3) phase at the composition of Bi8O8Se3Cl2 produced several mixed phases, as revealed by the XRD patterns. Most of XRD reflections can be indexed with a unit cell of a = 3.904(1) Å and c = 54.1580(1) Å with space group I4/mmm as indicated by the Pawley fit (FigureS2a), suggesting the existence of an (m = 1, n = 3) type structure. The dominant impurity peaks, marked with arrows, are ascribed to the (m = 1, n = 2) secondary phase, which is confirmed by significantly improved refinement (Rwp = 4.41%) of the XRD data after including the Bi6O6Se2Cl2 phase in Rietveld mode (FigureS2b). Furthermore, according to the stoichiometry, a second impurity phase with composition Bi2O2Se can be present, which was also confirmed by further fit improvement (Rwp = 4.08%) of the XRD data after considering Bi2O2Se in the Rietveld refinement (FigureS2c). Further evidence of the phase composition of the Bi8O8Se3Cl2 sample is provided in the following STEM-HAADF/ABF imaging section. Synthesis of the phase-pure sample of (m = 1, n = 3)-type Bi8O8Se3Cl2 was not successful although multiple reactions were attempted through changing the firing temperature and time as well as the stoichiometries.Crystal Structure of Bi6O6Se2Cl2. An ab initio method coupling was employed to determine the crystal structure of Bi6O6Se2Cl2 from powder diffraction data. First the XRD pattern of Bi6O6Se2Cl2 could be indexed with a tetragonal cell using the a = 3.8996 (1) Å and c = 19.6496(4) Å cell parameters. Systematic absences of reflections on XRD data correspond to h00, h = 2n+1 and hk0, h + k = 2n+1. These conditions suggest two probable space groups, P4/n or P4/nmm. Electron diffraction was then performed to confirm the symmetry and lattice parameters and gain more information about the Laue class. Selected-area electron diffraction (SAED) patterns recorded along the [100] and [001] zone axes (Figures5b and 5c) can be indexed with the tetragonal lattice parameters previously proposed by XRD. Then, the convergent beam electron diffraction (CBED) pattern recorded along the [001] zone axis contains a 4-fold axis and two sets of mirror planes (Figure5a), which suggest P4/nmm as the proper space group.

Figure 5 .

 5 Figure 5. (a) CBED pattern along [001] direction, and SAED patterns along (b) [100] and (c) [001]

Figure 6 .

 6 Figure 6. The Rietveld plots of (a) XRD and (b) NPD data for Bi6O6Se2Cl2 sample. The inset in (a)

  monolayer alternatively, which contains a central symmetric Bi2O2 2+ layer with bridging Se 2-layers on both sides, and two asymmetric Bi2O2 2+ layers adjacent to the terminal Cl -layers. The central oxygen layer of the latter shifts backward against the terminal Cl layer. Similar observations can also be made in the Bi4O4SeCl2 slab. Three distinct square antiprismatic Bi environments are present in the Bi6O6Se2Cl2 slab, all of which are coordinated to four oxide and four non-oxide anions. Due to the different nature of the Se layer and Cl layers, the bond lengths at the three Bi sites are different (Figure 7b). The Bi1 atom adjacent to the terminal Cl layer exhibits the shortest Bi-Cl/Se bonds of 3.040(2) Å, compensated with the largest Bi-O bonds of 2.376(3) Å. The Bi2 and Bi3 atoms adjacent to the bridging Se layer have longer Bi-Se/Cl bonds of 3.371(2) Å and 3.228(2) Å, and reduced Bi-O bonds of 2.267(3) Å and 2.320(1) Å, respectively.

Figure 7 .

 7 Figure 7. Multiple anion bonding environments form a composite charge-neutral slab in Bi6O6Se2Cl2. (a)

  Bi2O2Se and Bi2O2X2 layers elucidated above, we can reasonably suppose that the bridging Se layers in these intergrowth series has an ideal composition of 50% Se +50% Cl, which has the same Se defects content as that in the surface layer of Bi2O2Se if the single-phase Bi2+2nO2+2nSenCl2 compound can be stabilized. Then, terminal Cl layer of Bi2+2nO2+2nSenCl2 should have a Se: Cl ratio of n: 4-n based on the stoichiometry. Therefore, the terminal Cl layer in hypothetical stable single-phase Bi8O8Se3Cl2 (n = 3) and Bi10O10Se4Cl2 (n = 4) has a composition of 75% Se + 25% Cl and 100% Se, respectively. These Serich chemical compositions are very unlikely to be stable as a terminal layer due to the dominant presence of the Se elements, which prefers a 6-8 coordination with Bi forming a bridging Bi2O2Se-type interface rather than a 4-coordination in the terminal Bi2O2Cl2-type layer. The instability of the Bi8O8Se3Cl2 phase has been evidenced by the XRD and the following STEM-HAADF/ABF images, showing coexistence of different intergrowth structures or multiple phases, which include Bi8O8Se3Cl2 type structure, Bi6O6Se2Cl2 type structure, and Bi2O2Se phases as the dominated phases.The hypothetical (m = 2, n = 1) Bi6O6SeCl4 phase is expected to contain two successive vdW layered Bi2O2Cl2 slabs alternating with Bi2O2Se ionic layers, as shown in FigureS4, i.e. 2:1 superlattice of Bi2O2Cl2 and Bi2O2Se structural units. The unstable (m = 2, n = 1) structure may be also explained from the anionic disorder point of view. As indicated by the anionic mixing in the vdW layers of Bi4O4SeCl2 and Bi6O6Se2Cl2 containing ~ 30-40% Se, the Se content in the (m = 2, n = 1) structure could be not sufficient to maintain such an amount of anionic mixing in the high-order Bi2O2Cl2 slabs, which essentially destabilizes the supposed Bi6O6SeCl4 phase and transforms it into a mixed phase of Bi4O4SeCl2 and Bi2O2Cl2. The inaccessible high-order Bi2O2Cl2 slabs in the 2:1 superstructure are therefore in great contrast with the ionic high-order Bi2O2Se slabs showing variable thickness.STEM-HAADF/ABF imaging of Bi8O8Se3Cl2.As stated above, an (m = 1, n = 3) type structure has been identified in the Bi8O8Se3Cl2 sample. However, the Rietveld fit of the related XRD data using the (m = 1, n = 3) type structure is rather poor with several unfitted peaks (FigureS3).

Figure 8a and b show

  Figure 8a and b show a SAED pattern and a HRTEM image of the Bi8O8Se3Cl2 sample along the [100]

  middle): The bold dark column on the low-magnification STEM-HAADF image (Figure 8d) is confirmed as the double Cl layers, whilst the thin dark column corresponds to the single Se layers. As shown in Figure 8d, for most cases, the bold dark columns (double Cl layers) are well separated from each other at a distance > 2.0 nm by a thin dark column (Se layer), revealing the succession of multiple different stacking members, Bi6O6Se2Cl2, Bi8O8Se3Cl2 and other (m = 1, n >3) structures in one crystalline piece. Given that single-layer Bi2O2Cl2 and Bi2O2Se slabs have constant separations of ~ 0.8 nm and ~0.6 nm respectively, the layer thickness d of (m = 1, n) type intergrowth structure is related to the number of single-layer Bi2O2Se slabs n according to the relation d = 0.6n + 0.8 (nm). Therefore, based on the measured thickness d from the neighboring dark thick columns (double Cl layers), the number of Bi2O2Se slabs n in the (m = 1, n) type intergrowth structure can be derived, changing from bottom to top according to the sequence of 3 -3 -7 -2 -0 -2 -4 -6 -3 -5 -3 in zone 1. The most-populated intergrowth structure is the (m = 1, n = 3) type Bi8O8Se3Cl2 and the second one is the (m = 1, n = 2) type Bi6O6Se2Cl2 on zone 1 as shown in Figure 8d. In zone 2, as shown in Figure 8f, much larger interlayer distances (~ 5.0 nm -14.0 nm -8nm) are observed, which correspond to the (m = 1, n) type intergrowth structure with n values varying according to 6 -21 -11. This intergrowth structure has such a long c axis that the Bi2O2Cl2 layer appears almost as a planar defect in the Bi2O2Se lattice. The corresponding SAED pattern (Figure8a) has principal bright spots, which can be also indexed with the lattice parameters of Bi2O2Se, and many less intense spots between the main reflections, indicating the formation of complex superstructures due to the insertion of Bi2O2Cl2 layers. The materials are sensitive to electron beam (Figure8d) and will transform to a trigonal or hexagonal phase, which is probably the 6-coordinated anti-CaAl2Si2 structure (FigureS3a) under long-time electron radiation.

Figure 8 .

 8 Figure 8. On all the images the sample is oriented along the [100]* direction. (a) Typical SAED pattern

Figure 9 .Figure 11 .

 911 Figure 9. (a) UV/Vis diffuse reflectance spectra and (b) (F(R)hv) 1/2 vs hv plot of Bi2+2nO2+2nSenX2 samples.

Figure 10 .

 10 Figure 10. Band structures of (a) Bi6O6Se2Cl2 and (b) Bi4O4SeBr2 by DFT calculations. The orbital

Figure 11 .

 11 Figure 11. Charge densities of the VBM and CBM of (a) Bi4O4SeCl2, (b) Bi4O4SeBr2, (c) Bi6O6Se2Cl2,

Figure 12 .

 12 Figure 12. (a) Typical complex impedance plots for Bi4O4SeBr2 composition at 7 K and 240 K (inset).

Figure 13 .

 13 Figure 13. Imaging and orientation of the Bi4O4SeBr2 pellet. (a) SEM image, (b) experimental XRD

Figure 14 .

 14 Figure 14. Thermal conductivity versus temperature plots for Bi4O4SeX2 and Bi6O6Se2X2 (X = Cl and

1 )

 1 Temperature (K)of the materials show a similar optical band indirect gap of ~ 1.05 eV, which is related to the CBM and VBM originated by the electronic state in the Bi2O2Se slabs. The vdW materials possess low bulk resistivity (100 Ω cm) and exceptionally low thermal conductivity (0.66 W/mK and lower) parallel to the axis of pressing of the pellets, showing the heavier halogen, the lower thermal conductivity. The new multiple anion chemistry of layer stacking in the vdW materials complemented here will be helpful in the design of new vdW materials and their physical property tailoring.ACKNOWLEDGMENTThe National Science Foundation of China (No. 22090043, 51662013, 22161014, 21850410458), Guangxi Natural Science Foundation (No. 2019GXNSFGA245006, 2018AD19200, AD19245097, 2020GXNSFAA297220), and the Foundation of Guilin University of Technology (No. GUTQDJJ2018115) are acknowledged for the financial support. This project has benefited from the facilities of the Platform MACLE-CVL which was co-funded by the European Union and Centre-Val de Loire Region (FEDER).

Table 1 .

 1 Final refined crystallographic parameters and selected bond lengths for Bi4O4SeCl2*.

	Atoms	Site	x	y	z	Occupancy	Biso (Å 2 )
	Bi1	4e	0	0	0.56852(3) 1	0.03(2)
	Bi2	4e	0.5 0.5 0.66137(3) 1	0.03(2)
	O1	8g	0.5 0	0.61213(7) 1	0.03(5)
	Se1/Cl1 2a	0.5 0.5 0.5	0.45(1)/0.55(1)	0.45(6)

Table 2 .

 2 Final refined crystallographic parameters for Bi4O4SeBr2*.

	atoms	site x	y	z	occupancy	Biso (Å 2 )
	Bi1	4e	0	0	0.56819(3) 1	0.3(1)
	Bi2	4e	0.5 0.5 0.65656(3) 1	1.2(1)
	O1	8g	0.5 0	0.60895(9) 1	0.06(9)
	Se1/ Br1	2a	0.5 0.5 0.5	0.59(4)/0.41(4) 0.02(5)
	Br2/Se2	4e	0	0	0.70662(5) 0.79(2)/0.21(2) 0.02(3)
	Bond lengths (Å)
	Bi1-O1 ( 4)			2.276(2)	Bi2-O1 ( 4)	2.379(2)
	Bi1-Se1(Br1) ( 4)		3.378(1)	Bi2-Br2(Se2) ( 4)	3.116(1)
	* space group I 4/mmm, a = 3.93007(5) Å, c = 28.1763(6) Å, V = 435.196(9) Å 3 , ρ = 8.6899(2) g/cm 3 ,
				Rwp = 6.32%, χ 2 = 1.74.

Compared with its Cl analogue, the Bi4O4SeBr2 has a similar lattice parameter a but significantly longer c lattice parameter (by 1.147 Å) and thus larger volume. This variation is related to the larger ionic size of Br -(1.96 Å) than Cl -(1.81 Å)

Table 3 .

 3 Final refined crystallographic parameters of Bi6O6Se2Cl2*.

	(a.u.)		31.6	32.0	32.4	32.8	(a.u.)
	Intensity						Intensity
	20	40	60 2degrees 80	100	120	0.5	1.0	1.5 d-spacing 2.0	2.5 (Å)	3.0	3.5
	atom	site	x	y	z		occupancy	Biso (Å 2 )
	Bi1		2c	0.75 0.75	0.37812(7)	1		1.0(2)
	Bi2		2c	0.25 0.25	0.24950(7)	1		0.6(2)
	Bi3		2c	0.25 0.25	0.06436(8)	1		0.4(2)
	O1		2a	0.75 0.25	0		1		0.03(13)
	O2		4f	0.75 0.25	0.3087(3)	1		0.03(13)
	Se1/Cl1 2c	0.75 0.75	0.1503(2)	0.57(1)/0.43(1) 0.4(1)
	Cl2/Se2 2c	0.75 0.75	0.5561(2)	0.57(1)/0.43(1) 0.03(13)
	Bond lengths (Å)
	Bi1-O2 ( 4)		2.376(3)	Bi2-Se1/Cl1 ( 4)	3.371(2)
	Bi1-Cl2/Se2 ( 4)	3.040(2)	Bi3-O1 ( 4)		2.320(1)
	Bi2-O2 ( 4)		2.267(3)		Bi3-Se1/Cl1 ( 4)	3.228(2)

* space group P4/nmm, a = 3.8927 (1) Å, c = 19.621(1) Å, V = 297.32(2) Å 3 , ρ = 8.8173(2) g/cm

[START_REF] Zhou | Promising Thermoelectric Bulk Materials with 2D Structures[END_REF]

, Rwp = 7.16%, χ 2 = 2.6.

Supporting Information

The Supporting Information is available free of charge on the ACS Publications website at DOI: Comparison of the theoretical calculated and experimental measured lattice constants; The XRD data taken on the synthesis attempt of (m = 2, n = 1) phase Bi6O6SeCl4; Fits to the XRD plot for a synthesis

Notes

The authors declare no competing financial interests. CSD 2127176, 2126165, 2127172, and 2127058 contain the supplementary crystallographic data for this paper. The data can be obtained free of charge via www.ccdc.cam.ac.uk/data_request/cif, or by emailing data-request@ccdc.cam.ac.uk, or by contacting The Cambridge Crystallographic Data Centre, 12 Union Road, Cambridge CB2 1EZ, U.K.; fax: + 44 1223 336033.


Lijia Zhou #

Jinxiao Zhang

Congling Yin

Luhui Meng

Jiachen Li

Yuexiang Li

Franck Fayon

Mathieu Allix

Xiaojun Kuang
email: kuangxj@glut.edu.cn

Theoretical and Experimental Studies of Gallate Melilite Electrides from Topotactic Reduction of Interstitial Oxide Ion Conductor

Keywords: Electride, Nonstoichiometric Melilite, Interstitial oxide ions, Topotactic reduction

Non-stoichiometric La1.5Sr0.5Ga3O7.25 melilite oxide ion conductor features active interstitial oxygen defects in the pentagonal rings with high mobility. In this study, electron localization function calculation by density functional theory indicated that the interstitial oxide ions located in the pentagonal rings of the gallate melilites may be removed and replaced by electron anions, which are confined within the pentagonal rings, which would therefore convert the melilite interstitial oxide ion conductor into a zero-dimensional (0D) electride. The more active interstitials oxide ions, compared to the framework oxide ions, make the La1.5Sr0.5Ga3O7.25 melilite structure more reducible by CaH2 using topotactic reduction, in contrast with the hardly reducible nature of the parent LaSrGa3O7. The topotactic reduction enhances the bulk electronic conduction (σ ~ 0.003 S/cm at 400℃) by ~ one order of magnitude for La1.5Sr0.5Ga3O7.25. The oxygen loss in the melilite structure was verified and most likely takes place on the active interstitial oxide ions. The identified confinement space for electronic anions in melilite interstitial oxide ion conductors presented here provides a strategy to access inorganic electrides from interstitial oxide ion conductor electrolytes.

Introduction

Electrides that trap anionic electrons are an important class of ionic compounds demonstrating interesting chemical and physical properties such as strong reduction, metallicity, superconductivity, and catalytic activity owing to their unique electronic structures showing various anionic electron interconnectivity dimensionality from zero to three. [START_REF] Zhang | Recent Advances and Applications of Inorganic Electrides[END_REF][START_REF] Liu | Electrides: a review[END_REF] Electrides are classified as inorganic and organic types depending on the constituting elements of the framework structures. Most of electrides suffer from poor stability at room temperature and in air, therefore the synthesis of stable electrides is still challenging. Owing to stability advantages over organic electrides, the synthesis of inorganic electrides and the exploration of their potential applications have received considerable attentions. [START_REF] Dye | Electrons as Anions[END_REF] So far only a few inorganic electrides have been successfully prepared although theoretical studies predicted many more potential inorganic electrides. [START_REF] Srdanov | Na4 3+ ) clusters in sodium sodalite[END_REF][START_REF] Ichimura | Toward inorganic electrides[END_REF][START_REF] Wernette | Inorganic electrides formed by alkali metal addition to pure silica zeolites[END_REF] The well-known examples of inorganic electrides includes calcium aluminate mayenite Ca12Al14O33-based electrides (abbreviated as C12A7:e  )_ENREF_6 and dicalcium nitride Ca2N compound_ENREF_7 with various confinement spacings for anionic electrons covering from isolated to layered dimensionalities. [START_REF] Matsuishi | High-density electron anions in a nanoporous single crystal: [Ca24Al28O64] 4+ (4e[END_REF][START_REF] Lee | Dicalcium nitride as a two-dimensional electride with an anionic electron layer[END_REF] Among the known inorganic electrides, C12A7:e  electrides have been standing at frontier of inorganic electrides owing to their outstanding thermal and chemical stability at room temperature in air and intriguing functions for various applications. 1 _ENREF_1 [START_REF] Kitano | Ammonia synthesis using a stable electride as an electron donor and reversible hydrogen store[END_REF][START_REF] Toda | Activation and splitting of carbon dioxide on the surface of an inorganic electride material[END_REF][START_REF] Toda | Intense thermal field electron emission from room-temperature stable electride[END_REF][START_REF] Kim | Photoelectron Spectroscopic Study of C12A7:eand Alq3 Interface: The Formation of a Low Electron-Injection Barrier[END_REF][START_REF] Miyakawa | Superconductivity In An Inorganic Electride 12CaO•7Al2O3:e[END_REF] Ca12Al14O33 adopts a 3-dimensional (3D) tetrahedral network structure with subnanometer-sized cages which are partially occupied by extra oxygen atoms. The encaged oxygen atoms can be replaced by electron anions through topotactic reduction, leading to C12A7:e  electrides with variable concentration of anionic electrons and therefore displaying different colors. [START_REF] Matsuishi | High-density electron anions in a nanoporous single crystal: [Ca24Al28O64] 4+ (4e[END_REF] These electron anions are confined within the isolated cages, thus C12A7:e  represents a zero dimensional (0D) electride which may work as electron field emitter and cathode materials for organic light emitting diodes. [START_REF] Toda | Intense thermal field electron emission from room-temperature stable electride[END_REF][START_REF] Kim | Photoelectron Spectroscopic Study of C12A7:eand Alq3 Interface: The Formation of a Low Electron-Injection Barrier[END_REF] Initially, the anionic electrons in C12A7:e  was found to form diamagnetism [START_REF] Matsuishi | High-density electron anions in a nanoporous single crystal: [Ca24Al28O64] 4+ (4e[END_REF] and later to display superconductivity with very low critical temperature 0.16-0.4 K, being the first example of electride superconductor. [START_REF] Miyakawa | Superconductivity In An Inorganic Electride 12CaO•7Al2O3:e[END_REF] The electron anions in C12A7:e  can be replaced by hydride anions and the reversible storage and release of hydrogen in the cage make C12A7:e  an efficient catalyst for ammonia synthesis at low temperature and ambient pressure when combined with the conventional Ru catalyst. [START_REF] Kitano | Ammonia synthesis using a stable electride as an electron donor and reversible hydrogen store[END_REF][START_REF] Hayashi | Light-induced conversion of an insulating refractory oxide into a persistent electronic conductor[END_REF] However the recent investigation by Weber et al. concluded that the industrial application potential of C12A7:e  as catalyst for the ammonia synthesis is limited due to the irreversible hydride formation at pressures > 1 MPa. [START_REF] Weber | Mayenite-Based Electride C12A7e-: A Reactivity and Stability Study[END_REF]

_ENREF_7

As the topology of the confining cavity for anionic electrons has significant impact to the physical properties of electrides, it is important to identify new confinement spaces with unique topologies for anionic electrons in order to discover electrides with advanced performances. [START_REF] Lee | Dicalcium nitride as a two-dimensional electride with an anionic electron layer[END_REF] The discovery of the C12A7:e  electride by Hosono et. al. [START_REF] Matsuishi | High-density electron anions in a nanoporous single crystal: [Ca24Al28O64] 4+ (4e[END_REF] provides a useful structural route to access new inorganic electrides from the known structure prototypes of oxides containing extra oxygen atoms, which are not bonded to the polyhedral framework, as well as the hardly reducible cations e.g. alkaline (earth) and aluminum cations through topotactic reduction. This strategy indeed applies to the recently theoretical-predicted La8Sr2(SiO4)6:4e  silicate apatite 1-dimensional (1D) electride, [START_REF] Zhang | Electron Confinement in Channel Spaces for One-Dimensional Electride[END_REF] which was synthesized partially through removing the 1D oxygen channels bonding with the La/Sr cations only.

In the Ca12Al14O33 mayenite structure, the encaged extra oxygen atoms display mobility at high temperatures leading to oxide ion conductivity properties comparable to those of yttrium-stabilized zirconia YSZ. 17 _ENREF_15 The possibility to remove the mobile encaged oxygen atoms in the mayenite structure using topotactic reduction inspired our interest in interstitial oxide ion conducting electrolytes for discovery of new anionic electron confinement spacing and therefore oxide-based inorganic electrides. Indeed the environment for interstitial oxide ions could provide confinement space for electrons if these interstitial oxide ions can be removed. The LaSrGa3O7 Melilite structure is composed of alternative large cationic (La/Sr) layers and corner-sharing Ga3O7 tetrahedral layers (Figure 1a), featuring five-fold tunnels consisting of five corner-sharing GaO4 tetrahedra. The large (La/Sr) cations are eight-coordinated and located in the five-fold tunnels (Figure 1b).The gallate melilites have shown the ability to accommodate interstitial oxide ion defects with high mobility, which are incorporated into the non-planar five-fold rings and between two A cations (Figure 1c), transforming one of the tetrahedral GaO4 with terminal oxygen (O2) in the ring into a trigonal bipyramidal GaO5 (Figure 1d). [START_REF] Kuang | Interstitial oxide ion conductivity in the layered tetrahedral network melilite structure[END_REF] Recently, density function theory (DFT) calculations on the La1.5Sr0.5Ga3O7.25 composition revealed that the interstitial oxide ions have cooperatively weak antibonding interaction with framework oxygen atoms [START_REF] Xu | Chemical Bonding Effect on the Incorporation and Conduction of Interstitial Oxide Ions in Gallate Melilites[END_REF] , which make interstitials oxide ions highly active, i.e. they are loosely bound with the framework, which could be favorable for chemical reduction. Herein we demonstrate in principle that interstitial oxide ions in the pentagonal rings of the gallate melilites may be removed and replaced by electron anions which can be confined within the pentagonal rings and convert the melilite interstitial oxide ion conductors into 0D electrides and report its experimental verification. This work unravels a new cage-shaped confining space (Figure 1c) consisting of pentagonal rings and two A cations for the electron anions.

Methods

Synthesis. The La1+xSr1-xGa3O7+0.5x (x = 0, 0.5) polycrystalline samples prior to the topotactic reduction were prepared by hightemperature solidstate reaction using La2O3 (99.99%, Aladdin), Ga2O3 (99.9%, Aladdin), and SrCO3 (99.9%, Aladdin) as starting materials, which were weighed according to the correct stoichiometries with excess Ga2O3 (~ 3 mol%) and mixed thoroughly in ethanol using an agate mortar and pestle. The mixtures were preheated at 1200 ℃ for 12 h in air, then ground, uniaxially pressed into pellets and finally fired at 1400 ℃ for 12 h, leading to dense pellets with ~ 92% of X-ray theoretical densities.

The topotactic reduction of the La1+xSr1-xGa3O7+0.5x (x = 0, 0.5) samples was performed using CaH2 (Aladdin, > 98.5%) as a solid-state reducing agent. The as-made thin dense bulk ceramic samples (~ 1.4 mm thickness) from the high temperature sintering in air described above were buried in CaH2 powder (according to the mass ratio of melilite to CaH2, 0.5:1) in an argon-filled glovebox (O2 and H2O < 1 ppm), which were sealed in evacuated silica tubes and heated at various temperatures within 300-580 ℃ for 3-7 days for selecting the optimal temperature range for the topotactic reduction. The resulted pellets were washed with methanol solution of NH4Cl (0.1 M) to remove the calciumcontaining phases from ceramic surface (e.g., CaO and CaH2) and dried under vacuum at room temperature. The high densities were retained for the reduced ceramic pellets, which were polished to ensure further removal of surface impurities, and finally were used in the characterizations including the AC impedance spectroscopy measurements Characterization. The phase formation in the samples was investigated by X-ray powder diffraction (XRD), which was performed on a Panalytical X'Pert Pro diffractometer (Cu Kα radiation) equipped with an Anton Paar oven chamber (HTK1600N). High-quality XRD data suitable for Rietveld refinement was collected over the 2θ range of 10-120˚ at room temperature. Variable temperature X-ray powder diffraction (VT-XRD) data were collected under air at room temperature and different temperatures with a 50 ℃ temperature step from 50 ℃ up to 1000 ℃ over the 2θ range from 10° to 80°. Neutron powder diffraction (NPD) data with a constant wavelength (λ = 1.225 Å) were collected at room temperature on the 3T2 powder diffractometer at the Laboratoire Léon Brilloin (Saclay, France) over a 2θ range of 5-120°. Rietveld refinements were performed on both XRD and NPD data using Topas Academic 6.0 [START_REF] Coelho | TOPAS and TOPAS-Academic: an optimization program integrating computer algebra and crystallographic objects written in C plus[END_REF] with the modified Thompson-Cox-Hastings pseudo-Voigt function for the profile refinement. [START_REF] Thompson | Rietveld refinement of Debye-Scherrer synchrotron X-ray data from Al2O3[END_REF][START_REF] Young | The rietveld method[END_REF] 1 H solid-state nuclear magnetic resonance (NMR) experiments were performed on a Bruker Neo spectrometer operating at a magnetic field of 20.0 T corresponding to a 1 H Lamor frequency of 850 MHz. [START_REF] Zhang | Recent Advances and Applications of Inorganic Electrides[END_REF] H quantitative magic angle spinning (MAS) spectra were recorded at a spinning frequency of 30 kHz using a Hahn echo sequence. The echo delay was set to one rotor period (33.33 µs) and 32 transients were collected using a recycle delay of 3 s. The 1 H NMR spectra of the materials were obtained after subtracting the proton signal of an empty rotor recorded under the same experimental conditions.

The oxidation thermogravimetric analysis (TGA) was performed under atmosphere of O2 at a flowing rate of 20 ml/min using a STA8000 thermal analyzer from 30 ℃ to 1000 ℃ to determine the oxygen content in the topotactic reduced samples. Magnetization measurements were performed on a Quantum Design MPMS magnetometer under a 1000 Oe field while the samples were warmed from 2 K to 300 K, following zerofield cooling (ZFC) and field cooling (FC). AC Impedance spectroscopy (IS) measurements were carried out with a Solartron 1260A frequency response analyzer over the 10 - 1 10 6 Hz frequency range on the pellet samples placed between two silver electrodes on heating from 25 ℃ to 400 ℃ in a 5% H2-95% N2 atmosphere. X-ray photoelectron spectroscopy (XPS) measurements were carried out on an ESCALAB 250Xi spectrometer using Al Ka X-ray as the excitation source.

The optical diffuse reflectance data were collected in the 200-800 nm range using a Shimadzu (Kyoto, Japan) UV-3600 spectrometer equipped with an integrating sphere. Water reduction activity was examined under ultraviolet-visible (UV-Vis) light irradiation with a Xe lamp (350 W) or under dark in closed Pyrex cells. Typically, 0.5 g sample and 100 mL of the formic acid (0.1 M) containing chloroplatinic acid (2 wt% Pt) were added into the cell. Before irradiation, the mixture was treated with ultrasound for 10 min, and then aerated with N2 bubbling for 25 min. The amount product H2 was determined on a GC 1690-type gas chromatograph using N2 gas as carrier for H2 detection.

DFT calculations. The formation energies and electronic structures of interstitial-containing

La1.5Sr0.5Ga3O7.25, as well as the reduced compositions La1.5Sr0.5Ga3O7 and La1.5Sr0.5Ga3O7.125 were investigated by density functional theory using Vienna Ab Initio Simulation Package (VASP, version 5.4.1). [START_REF] Kresse | Efficient Iterative Schemes for Ab Initio Total-Energy Calculations Using a Plane-Wave Basis Set[END_REF] Generalized gradient approximation (GGA) with Perdew-Burke-Ernzerhof functional (PBE) [START_REF] Perdew | Generalized Gradient Approximation Made Simple[END_REF] and projector augmented wave (PAW) method were adopted for all the simulations. The plane-wave cutoff energy was set to 500 eV and a Γ-point of 4×4×3 was used for geometry optimization and 6×6×5 for electronic structure analysis to sample the first Brillouin zone. Geometry optimizations without structural constraints were carried out with convergence criterions for energy and force at 1.0 × 10 -5 eV/atom and 0.02 eV/Å, respectively. A 1×1×2 supercell (see models in Figure 2, a = 8.05 Å, c = 10.66 Å) with 49 atoms for La6Sr2Ga12O29 with interstitial oxide defects or 48 atoms for La6Sr2Ga12O28 without interstitial oxide defects and a 2×1×2 supercell was created by doubling the unit cell along the a-axis direction to realize the partial removal of interstitial oxide in composition La12Sr4Ga24O57 (La1.5Sr0.5Ga3O7.125). The valence states of the structures include 5s 2 5p 6 5d 1 6s 2 for La atoms, 4s 2 4p 6 5s 2 for Sr atoms, 3d 10 4s 2 4p 1 for Ga atoms and 2s 2 2p 4 for O atoms.

The interstitial oxide ion defect formation energies can be defined according to equation ( 1) and ( 2

): 25 La4Sr4Ga12O28(s) + La2O3(s) → La6Sr2Ga12O29(s) + 2SrO(s) (1) △Hf = [E(La6Sr2Ga12O29(s) + 2E(SrO(s))] -[E(La4Sr4Ga12O28(s)) + E(La2O3(s))] (2)
Where E(X) is the total energy of the corresponding species.

The typically theoretical analyses for electrides including the density of states, band structures, 3D charge distribution and 2D electron localization function were performed with DFT calculations for the melilite compositions. [START_REF] Hosono | Advances in materials and applications of inorganic electrides[END_REF][START_REF] Li | Electron-Deficient-Type Electride Ca5Pb3: Extension of Electride Chemical Space[END_REF][START_REF] Li | Discovery of Electrides in Electron-Rich Non-Electride Materials via Energy Modification of Interstitial Electrons[END_REF][START_REF] Dale | Theoretical Descriptors of Electrides[END_REF][START_REF] Wu | Intermetallic electride catalyst as a platform for ammonia synthesis[END_REF] More details on the DFT calculations are provided in the supporting information.

Results and discussion

First, DFT calculations were performed on the La1. Electron localization function (ELF) calculated with DFT method can be a reliable theoretical descriptor for us to evaluate whether the electrons left by the escaped interstitial oxide ions are caged as anions in the reduced La1.5Sr0.5Ga3O7, i.e. La1.5Sr0.5Ga3O7 forms as electride. [START_REF] Dale | Theoretical Descriptors of Electrides[END_REF][START_REF] Li | Is mayenite without clathrated oxygen an inorganic electride?[END_REF] The values of ELF are at the range of 0 ~ 1: 0 indicates that the electrons are highly delocalized or there is no electron, while 1 indicates totally localized electrons. Figure 4 a and e show the optimized structures of La1.5Sr0.5Ga3O7.25 and La1.5Sr0.5Ga3O7, respectively. The three-dimensional (3D) plot of ELF distribution results reveals that in La1.5Sr0.5Ga3O7 the electrons are highly localized at the position of interstitial oxygen after its removal (Figure 4 b, f), i.e., the electrons left by the escaped interstitial oxygen atoms are indeed caged, the integrated extra caged electron is about 0.88 e -. In order to observe the caged electrons more easily, we plot two-dimensional (2D) maps of ELF distribution along (001) plane and the corresponding max ELF value is about 0.74, which shows that the interstitial electrons are located at the center of the pentagonal ring. Thus, the reduced La1.5Sr0.5Ga3O7 can be described as an electride and formulated as 4(La1.5Sr0.5Ga3O7):2e -. In order to produce an electride material from La1.5Sr0.5Ga3O7.25, topotactic reduction was performed on La1.5Sr0.5Ga3O7.25 containing interstitial oxygen at various temperatures ranging from similar XRD patterns with that for the as-made one (Figure 5a), except for partial decomposition due to over-reduction at 580 ℃, which resulted in a small amount of La2O3 and LaSrGaO4 phases (Figure 5a and S4a). With the increase of the reduction temperature, the color of La1.5Sr0.5Ga3O7.25 sample gradually changed from white to black, as show Figure 5b. In contrast with the interstitials-containing composition, the reduction of the parent LaSrGa3O7 without interstitial oxygen at 580 ℃, did not much change the color (Figure 5c) and phase (Figure S5) of the sample. The ease to reduce the interstitials-containing nonstoichiometric composition compared with the stoichiometric composition may therefore be ascribed to the existence of interstitial oxide ions. In other words, the interstitials oxide ions are more active than the framework oxide ions in the melilite structure. The refined cell parameters of the samples before and after reduction are almost identical except for the slightly smaller cell volume for the La1.5Sr0.5Ga3O7.25- sample reduced at 580 ℃ (Figure S6), on which the following characterizations were performed if not specified. In order to examine whether the H -hydride anion was incorporated into the melilite structure and the Ga valence change in La1.5Sr0.5Ga3O7.25 during the CaH2 reduction, 1 H MAS NMR and XPS measurements were carried out. As shown in Figure S7, [START_REF] Zhang | Recent Advances and Applications of Inorganic Electrides[END_REF] H MAS NMR measurements excluded the possibility of H -hydride anion in the melilite structure. [START_REF] Hayashi | Hydride ions in oxide hosts hidden by hydroxide ions[END_REF] The XPS data show apparent existence of metal Ga signal in the reduced La1.5Sr0.5Ga3O7.25- samples (Figure S8). The EDS element mapping for the reduced samples (Figure S9) show that there is aggregation of Ga at the surface of the ceramic sheet, confirming that the topotactic reduction partially reduces the melilite materials to produce metal Ga. It should be noted here that the high Ga metal content from the XPS data does not means that most of Ga 3+ in the samples was reduced to Ga metal as the XPS only probes the sample surface (only a thickness of a few nanometers). Moreover Ga metal is liquid at room temperature, therefore easily immersing the grains along the surface. The total Ga content in the reduced samples is rather limited as indicated by the TGA data and water reduction characterization described below.

The existence of Ga metal in the reduced La1.5Sr0.5Ga3O7.25- samples was also confirmed by the magnetism measurement, which show normal diamagnetic behavior above 10 K but significantly enhanced diamagnetic susceptibilities below 7.22 K (Figure 6 a-b). This behavior could be ascribed to the superconductivity of metal Ga in the reduced samples. However, the reduced samples were found to be highly resistive all through the low temperature region below room temperature, which confirms that the reduced samples had very low contents of Ga metal, which did not connect with each other in the reduced samples to produce the short circuit effect. On the other hand, the as-made sample displayed single 2p3/2 and 2p1/2 XPS peaks for Ga 2p orbits (Figure S8b) and diamagnetic behavior as normal, followed by a transition to paramagnetic substance below 7.22 K (Figure 6c).

Such magnetic behavior suggests absence of Ga metal in the as-made sample. Therefore, the XPS peaks for Ga 2p orbits all come from Ga 3+ instead of the mixed Ga and Ga 3+ found in the reduced samples. Thermal gravimetric analysis (TGA) was performed on the reduced La1.5Sr0.5Ga3O7.25- samples in O2, which showed a change in color from black to white and an associated weight increase ~ 0.51 wt% (Figure S10). This mass gain started from about 400 C, is ascribed to reoxidation of the reduced La1.5Sr0.5Ga3O7.25- sample, from which the removed oxygen content  per formula can be calculated as 0.18, assuming that the mass gain was completely from the reoxidation of reduced melilite phase.

The reoxidation of the sample was also confirmed by the VT-XRD experiment in air (Figure S11),

showing that the sample retained the tetragonal melilite structure after oxidation, except for the formation of LaGaO3, probably due to the volatilization of Ga2O3 at high temperatures. Although there was no apparent Ga2O3 phase observed in the VT-XRD data during the oxidation procedure, the oxidation of Ga metal to Ga2O3 can take place from ~ 400 C and the Ga2O3 could appear as amorphous. Apart from this, there could be still oxygen uptake for the melilite phase, i.e., the oxygen atoms in the melilite phase could be removed partially during the topotactic reduction and the reoxidation occurred for the melilite phase as well. Considering that the stoichiometric LaSrGa3O7 is hard to be reduced, the interstitial oxide ions are suspected to be removed during the topotactic reduction of in La1.5Sr0.5Ga3O7.25. Although the Ga metal content and the removed oxygen content in the melilite phase (which is lower than the determined  value per formula above) are not able to be discerned from the oxygen-oxidation TGA data, the electrons associated with the oxidation can be determined as 6.25  10 -4 electrons per 1g reduced sample.

The reduced La1.5Sr0.5Ga3O7.25- samples were found to react with water under dark environments, producing hydrogen gas (Figure 7). The hydrogen evolution can be enhanced by both visible and UV light radiation. However no apparent hydrogen evolution was observed at the visible and UV light radiation time exceeding 8 hours. This result indicates that the reduced samples do not have real photocatalytic activity and the hydrogen evolution could be mainly arising from the reactivity of metallic gallium with the water. The enhanced hydrogen evolution under light irradiation may be ascribed to the light absorbance of the sample accelerating the electron releasing from the metallic Ga and therefore the reduction reaction. The electrons associated with the saturated hydrogen evolution under light radiation (Figure 7) was estimated as 3.67  10 -4 electrons per 1g reduced sample. This value is smaller than that (6.25  10 -4 electrons per 1g reduced sample) associated with the oxidation determined from the TGA data above. The discrepancy between the electron numbers from different experiments could not be unmeaningful and the possibility that some electrons in the reduced sample did not participate the hydrogen evolution can not be excluded completely, which could be in the melilite phase. NPD data were collected on 6g reduced La1.5Sr0.5Ga3O7.25- sample, which was obtained through gathering the products from different reduction batches. Rietveld refinement of NPD data (Figure S12) was performed based on the average structural model of La1.54Sr0.46Ga3O7.27, [START_REF] Kuang | Interstitial oxide ion conductivity in the layered tetrahedral network melilite structure[END_REF] to analyze the oxygen content in the reduced product, which converged to Rwp ≈ 4.29% and Rp ≈ 3.36%. The occupancies for the oxygen sites were freely refined, which showed that the framework oxygen sites are essentially fully occupied (The refined occupancies were 0.98(1), 1.00(8) and 0.996 (8) for O1, O2 and O3, respectively and fixed as units) and the interstitial O4 site has an occupancy of 9.4(4)% (its refined atomic displacement factor Biso is 0.5(3) Å 2 , Table S1), apparently lower than that (12.5%) corresponding to the pristine La1.5Sr0.5Ga3O7.25 composition. [START_REF] Kuang | Interstitial oxide ion conductivity in the layered tetrahedral network melilite structure[END_REF] When the occupancy of O4 was fixed at 12.5% (corresponding to the full interstitial oxide ion content without reduction), its refined Biso increased to 1.45 Å 2 with the reliability factors slightly increased to Rwp ≈ 4.36% and Rp ≈ 3.42%.

Considering that the occupancy is highly correlated with the atomic displacement factor, we cannot determine precisely the interstitial oxygen loss based on the structural analysis on NPD data. It should be noted here that there are still some misfits in the NPD data fitting (Figure S12). Such misfit could be ascribed to the positional disorder of framework atoms arising from the incorporation of interstitial oxide ion, which is beyond description of the average structure model employed here. [START_REF] Kuang | Interstitial oxide ion conductivity in the layered tetrahedral network melilite structure[END_REF] In order to examine how the reduction affects the electrical conductivities of La1.5Sr0.5Ga3O7.25, AC impedance data of the reduced La1.5Sr0.5Ga3O7.25- pellet were collected on the as-made La1.5Sr0.5Ga3O7.25 pellet for comparison. Figure 8a shows the typical complex impedance plot of the as-made La1.5Sr0.5Ga3O7.25 at 150 ℃, which consists of a semicircular arc from bulk response with capacitances ~ 10 -12 -10 -11 F/cm and a Warburg-type electrode response with large capacitances ~ 10 - 8 10 -7 F/cm in the low frequency range, consistent with the oxide ion conduction. [START_REF] Irvine | Characterization by Impedance Spectroscopy[END_REF] There is intermediate arc between the bulk and electrode response arcs, showing ~ 10 -10 F/cm capacitances, indicative of grain boundary response. The bulk response arcs can be modeled with an equivalent circuit consisting of three parallel Rb, Cb, and (CPE)b elements (b is for bulk, and CPE is a constant phase element). [START_REF] Johnson | ZView: a software program for IES analysis, version 2.8[END_REF] The fitted Cb is ~ 7 pF/cm , consistent with the bulk response. While for the reduced La1.5Sr0.5Ga3O7.25- sample, the high-frequency arc showed capacitances covering from 10 -12 F/cm to 10 -10 F/cm at 150 ℃ indicating overlapped bulk and grain boundary response arcs, which cannot be accurately distinguished even by equivalent circuit modeling, and a bended Warburg-type electrode response arc (Figure 8b), indicative of mixed electronic and ionic conduction. Figure 8c shows the total conductivities of the reduced La1.5Sr0.5Ga3O7.25- sample, varying within 10 -7 -10 -3 S/cm from 100 ℃ to 400 ℃, which is ~ 1 order of magnitude higher than the bulk conductivities for the as-made La1.5Sr0.5Ga3O7.25. This result reveals that the bulk conductivities of reduced sample are higher than that for the as-made sample. Accordingly, the reduced sample displayed lower activation energy (Ea) ~ 0.74 eV than that (~ 0.82 eV) for the as-made materials.

Here the influence of phase decomposition on the electrical conduction behavior upon the reduction is considered to be limited based on the following three aspects. (i) The electrically insulating nature of the reduced samples above room temperature confirms the low content of Ga metal and the Ga metal liquid appeared isolated. Therefore, the metallic gallium certainly is not the case causing the enhanced bulk conductivity. (ii) Given the low contents (covering from ~ 0.7% to ~ 4% from NPD and XRD data) for LaSrGaO4 and La2O3 phases in the reduced samples and their much lower conductivity than the melilite oxide ion conductors, it is unlikely that the enhanced conductivity arose from these two phases. [START_REF] Raghvendra | Electrical properties of Ba doped LSGM for electrolyte material of solid oxide fuel cells[END_REF][START_REF] Wang | Tuning La2O3 to high ionic conductivity by Ni-doping[END_REF] _ENREF_38_ENREF_38_ENREF_38 (iii) The existence of lanthanum hydroxide was considered here as this phase was observed in the powder sample for the NPD experiment.

However, the pellets for the impedance measurement of the reduced sample were obtained through burying the as-made dense pellets in CaH2 powder. The dense nature of the pellets was retained after the reduction. The moisture in air is hardly to go through such dense pellets during the impedance measurement at high temperature. No apparent lanthanum hydroxide phase was observed in the XRD data of the pellet collected before and after the impedance measurement. Thus, neither the lanthanum hydroxide phase contributed to the enhanced conductivity. In short words, it is highly likely that the enhanced bulk conductivity upon reduction is due to the topotactic reduction of the melilite phase, rather than the decomposition. In other words, the oxygen loss in the melilite phase could indeed took place topotactically in the reduction of La1.5Sr0.5Ga3O7.25 although it is mild. Taking the mixed conduction observed from the impedance data into consideration, the enhanced bulk conductivity for the reduced sample comes from the contribution of electronic conductivity, which could arise from the electrons left by the escaped oxide ions in the melilite phase upon reduction. As indicated by the hardly reducible nature for the parent LaSrGa3O7 and the structural analysis on the NPD data described above, oxygen loss is highly likely on the active interstitial oxide ions rather than the framework oxide ions, given that the loss of framework oxide ions will break the corner-sharing connectivity for the tetrahedral layer, and the removing of interstitial oxygen atoms in La1.5Sr0.5Ga3O7.25 does not change the framework at all.

_ENREF_33 Considering that the interstitial oxides were not completely reduced during topotactic reduction, we have also constructed a 2×1×2 supercells to simulate La1.5Sr0.5Ga3O7.125 using configuration 1 to probe the electronic structure in the partially reduced melilite. This composition shared similar results on the electronic structure (Figure S13) and ELF value (Figure S14) with those of 1×1×2 supercell for La1.5Sr0.5Ga3O7. This means that when the individual interstitial oxide ion is removed, the electrons left by the escaped oxygen still act as anionic electrons, leading to nonstoichiometric electride.

Conclusion

DFT ELF calculations indicated that the interstitial oxide ions in the pentagonal rings of the gallate melilites may be removed and replaced by electron anions, which are confined within the pentagonal rings and convert the melilite interstitial oxide ion conductor into 0D electrides. In contrast with the hardly reducible nature for the parent LaSrGa3O7, the topotactic reduction of La1.5Sr0.5Ga3O7.25 by CaH2 resulted in black product with small amount of Ga 3+ reduced to Ga metal.

This reveals that interstitials oxide ions are more active than the framework oxide ions, making the melilite structure more reducible. The enhanced bulk electronic conduction by ~ one order of magnitude in the reduced samples verified the oxygen loss in the melilite structure during the topotactic reduction, which was highly likely on the active interstitial oxide ions. The new confinement space for electronic anions identified from the melilite interstitial oxide ion conductors in this work suggests that more interstitial oxide ion conductor electrolytes can be candidate hosts for preparation of new inorganic electrides.

Figure 1 .

 1 Figure 1. Crystal structure for interstitial oxide ion conducting gallate Melilite LaSrGa3O7. (a)

  5Sr0.5Ga3O7 composition (here the interstitial oxide ions are completely removed) with La1.5Sr0.5Ga3O7.25 for comparison. Initially, the favorable structures for DFT calculations were searched based on the interstitials-containing La1.5Sr0.5Ga3O7.25 composition, and transferred to the reduced compositions. When La atoms replace Sr atoms to generate interstitial oxide ions for neutralization, the interstitial oxide ions may appear in the pentagonal ring between two La, two Sr or one La and one Sr. Based on these considerations, we constructed 6 possibly representative configurations in 1 × 1 × 2 supercells which contain one interstitial oxygen for each supercell, as shown in Figure2, and calculated their formation energies (ΔHf) according to Eq. (2) through the geometry optimization. The calculations indicate that configurations 1 and 5 hold smaller formation energies compared with the other 4 configurations and configuration 5 has the lowest formation energy (Figure2). Both configurations 1 and 5 were simulated in this study. However we present the results for configuration 1 in the main text taking the following three aspects into consideration: (i) La 3+ shows higher positive charge than that of Sr 2+ , which results in stronger electrostatic interactions with interstitial oxide ; (ii) besides, La 3+ (r = 1.16 Å) holds smaller ionic radius compared with Sr 2+ (r = 1.26 Å), indicating the smaller steric hindrance for the incorporation of interstitial oxide ions[START_REF] Schutt | Oxygen Ion Migration and Conductivity in LaSrGa3O7 Melilites from First Principles[END_REF] (iii) there are more La-Oi-La configurations than La-Oi-Sr or Sr-Oi-Sr configurations in La1.5Sr0.5Ga3O7.25 from a statistical point of view. The theoretical simulation results of configurations 5 are provided in the supporting information (FigureS1 and S2).

Figure 2 .

 2 Figure 2. The 6 configurations in 1 × 1 × 2 supercells illustrating different distributions of La, Sr and

Figure 3

 3 Figure3shows the band structures, total density of states (TDOS) and partial/projected density of

Figure 3 .

 3 Figure 3. (a, c) Band structures and (b, d) total density of states (TDOS) and partial/projected density

Figure 4 .

 4 Figure 4. Optimized structures and computed ELF distributions of (a-d) La1.5Sr0.5Ga3O7.25 and (e-h)

Figure 5 .

 5 Figure 5. (a) XRD patterns and (b) color evolution of the La1.5Sr0.5Ga3O7.25- samples reduced at

Figure 6 .

 6 Figure 6. (a) c vs temperature plot for the reduced La1.5Sr0.5Ga3O7.25- recorded in the ZFC and FC

Figure 7 .

 7 Figure 7. Hydrogen evolution over the reduced La1.5Sr0.5Ga3O7.25- sample in the formic acid solution

Figure 8 .

 8 Figure 8. (a, b) The typical complex impedance plots at 150℃ and (c) bulk conductivities of (a) the

Acknowledgement

National Natural Science Foundation of China (No. 22090043, 21622101, 22103019), Guangxi Natural Science Foundation (No. 2019GXNSFGA245006), Guangxi Key Laboratory of Electrochemical and Magneto-chemical Functional Materials (No. EMFM20211111) are acknowledged for the financial support. The authors thank Florence PORCHER for neutron experiments realized at the Laboratoire Léon Brillouin in Saclay, France and Prof. Xing Ming for the valuable discussion on the DFT calculation results.

ASSOCIATED CONTENT

Supporting information

Comparison on DOS, optimized structures and computed ELF distributions of La1.5Sr0.5Ga3O7.25 and La1.5Sr0.5Ga3O7 based on configuration 5; Comparison of the TDOS between La1.5Sr0.5Ga3O7.25 and La1.5Sr0.5Ga3O7/La1.5Sr0.5Ga3O7.125 based on configurations 1 and 5; Comparison on DOS, optimized structures and computed ELF distributions of La1.5Sr0.5Ga3O7.25 and La1.5Sr0.5Ga3O7.125 based on configuration 1; XRD patterns and Ga 2p XPS spectra of the as-made and reduced LaSrGa3O7 samples; Rietveld plots for XRD data and the cell parameters of the reduced La1.5Sr0.5Ga3O7.25- and LaSrGa3O7 products; insitu variable temperature XRD patterns, 1 H MAS NMR spectra, Rietveld plot for NPD data, the final refined structural parameters, EDS elemental mapping and thermogravimetric data in oxygen atmosphere of the reduced La1.5Sr0.5Ga3O7.25- . This information is available free of charge at https://pubs.acs.org/.

Author Contributions

Lijia Zhou and Jinxiao Zhang contributed equally.

Notes

The authors declare no competing financial interest.

Table of Contents synopsis

The interstitial oxide ions located in the pentagonal rings of the non-stoichiometric La1.5Sr0.5Ga3O7.25 melilite oxide ion conductor may be removed and replaced by electron anions, which are confined within the pentagonal rings, and therefore convert the melilite interstitial oxide ion conductor into a zero-dimensional (0D) electride.