Computationally Guided Discovery of the Sulfide Li3AlS3 in the Li-Al-S Phase Field: Structure and Lithium Conductivity.

With the goal of finding new lithium solid electrolytes by a combined computational-experimental method, the exploration of the Li-Al-O-S phase field resulted in the discovery of a new sulfide Li3AlS3. The structure of the new phase was determined through an approach combining synchrotron X-ray and neutron diffraction with 6Li and 27Al magic-angle spinning nuclear magnetic resonance spectroscopy and revealed to be a highly ordered cationic polyhedral network within a sulfide anion hcp-type sublattice. The originality of the structure relies on the presence of Al2S6 repeating dimer units consisting of two edge-shared Al tetrahedra. We find that, in this structure type consisting of alternating tetrahedral layers with Li-only polyhedra layers, the formation of these dimers is constrained by the Al/S ratio of 1/3. Moreover, by comparing this structure to similar phases such as Li5AlS4 and Li4.4Al0.2Ge0.3S4 ((Al + Ge)/S = 1/4), we discovered that the AlS4 dimers not only influence atomic displacements and Li polyhedral distortions but also determine the overall Li polyhedral arrangement within the hcp lattice, leading to the presence of highly ordered vacancies in both the tetrahedral and Li-only layer. AC impedance measurements revealed a low lithium mobility, which is strongly impacted by the presence of ordered vacancies. Finally, a composition-structure-property relationship understanding was developed to explain the extent of lithium mobility in this structure type.

Introduction

All solid-state batteries (ASSBs) are of considerable current interest because they are a potential route to the use of lithium metal anodes while avoiding dendrite formation.[1,2] Solid-state electrolytes (SSEs) offer advantages over liquid electrolytes, such as large electrochemical stability windows and better thermal stability,[3,4] and have a lithium transport number of unity. Moreover, inorganic crystalline lithium ion conductors have superior ionic conductivity compared to organic polymers.[5] Recent reviews compared the different families of Li electrolytes,[5−9] and sulfides have the highest Li ion conductivity. For instance, members of the thio-LISICON family[10−14] such as the Li7P3S11 crystalline phase discovered in the Li2S–P2S5 system[15,16] and the Li–argyrodites Li6PS5X (X = Cl, Br, I)[17,18] have superior ionic conductivities, σ, ∼1 to 20mScm–1 at room temperature (RT) and low activation energies, Ea, ∼0.20 eV. These values are competitive with those of liquid electrolytes and make them promising candidates for integration in ASSBs. The sulfide anion is larger than oxide and more polarizable and affords lower frequency lattice vibrations, all favoring higher lithium mobility,[19−21] while superior mechanical properties offer lower grain boundary resistance[6] and allow easier processing into dense pellets. While these characteristics indicate superiority of sulfides over oxides, the usually narrower electrochemical stability window of the former can lead to parasitic reactions at the electrodes and considerably limit the performance and life cycle of ASSBs containing sulfide electrolytes.[5,11,22] Moreover, sulfide materials are often not stable in air, which makes them more difficult to handle. Further improvement of SSEs is needed for their commercialization, which implies finding materials showing good ionic conductivity as well as being able to form stable and conducting interfaces at the cathode and anode sides.[8,14,23]

Given the large number of candidate material compositions for potential new SSEs, computation has been extensively used to guide experimental work. On the one hand, some studies focus on currently known materials in order to produce an understanding of the underlying mechanisms for ion transport thanks to density functional theory (DFT) energy landscape methods and molecular dynamics.[20,24,25] The need for higher throughput screening methods led to database-driven approaches where models of ion transport, such as the development of predictive performance metrics,[26] are used to prioritize existing materials for evaluation as SSE. This can be done for example through the use of machine learning algorithms combined with first-principles calculations[27,28] or with bond valence mapping analysis.[18,29−31] On the other hand, the discovery of new materials with unique compositions and/or structures can be accelerated by predicting the most stable compounds within a chosen phase field prior to experimentation. The crystal structure for which the energy is calculated can be either thought to be isostructural to a known material in which original substitutions are performed[32−34] or can be determined through crystal structure prediction techniques.[35−37]

A relatively few mixed anion systems have been studied as SSE. Such materials can combine the advantages of each individual anion and may also offer new structure types for ion transport. The good performance of argyrodite sulfide halide materials demonstrates the potential of such mixed anion systems.[18,38] Oxysulfide phases, especially for lithium solid electrolyte applications, however remain less explored[39−42] and could be a promising approach to yield new materials. In particular, the oxygen for sulfur substitution in LGPS structures has been an area of investigation.[40,42] For instance, Kim et al.[42] reported the study of the Li10SiP2S12–O (0 < x < 1.75) solid solution, which resulted in an optimal conductivity for x = 0.7. Another interesting example is the theoretical study by Wang et al., presenting O doping into β-Li3PS4, which showed that the oxygen insertion enabled lower activation energy barriers, improved electrochemical stability, and created a beneficial 3D connection pathway for Li diffusion.[41] Moreover, among non-crystalline materials, oxysulfide glasses such as combinations of Li2S–P2S5–P2O5[43] or Li2S–SiS2–LiBO2[44] have been studied as potential lithium solid electrolytes and showed promising properties.

In this study, we evaluate oxysulfides of aluminum as candidate lithium ion conductors due to the earth abundance, low cost and toxicity, high polarizability, and redox inactivity. The composition LiAlOS was previously identified as a new potential interesting solid electrolyte target in a computation-only study.[37] The present study of the Li–Al–O–S phase diagram was guided by computational evaluation of the stability of candidate compositions, which both highlights where new phases are likely to be found and allows assessment of the amount of experimental effort to be invested at each composition. Within this phase diagram, the new sulfide phase Li3AlS3 was identified and its structure and lithium ion conductivity were characterized.

Experimental Section

Synthesis

Materials

Li2S (99.98%), LiNO3 (>99%), Li2CO3 (>99%), 7Li2CO3 (99% 7Li), and Al(OH)3 (reagent grade) were purchased from Sigma Aldrich, while Li2O (99.5%), urea (99.0%), Al2S3 (99%+), and Al(NO3)3·9H2O (98%) were obtained from Alfa Aesar.

Exploratory Synthesis of the Compounds in the Li–Al–O–S Phase Field

For all sulfide-containing materials, precursors and resulting powders were handled in a He-filled glovebox. Compositions belonging to the Li–Al–O–S phase diagram were synthesized from stoichiometric mixtures of Li2O, Al2S3, and, when needed, Li2S and pre-synthesized LiAlO2 (according to a procedure from Gao et al.,[45]cf. the Supporting Information, SI). The precursors were weighed in the appropriate amount in order to obtain a total mass of 300 mg. The powders were then mixed and ground in an agate mortar for 15 min, transferred to an alumina crucible, and placed in a quartz tube before sealing under vacuum (10–4 mbar). The tube containing the sample was heated to 800 °C at a ramp rate of 5 °C·min–1, held at 800 °C for 48 h, and then quenched in water. The resulting powder was then manually ground in order to obtain a fine powder.

Final Synthesis of Li3AlS3

After identifying Li3AlS3 as a new phase through the described synthesis method above, its synthesis was slightly modified in order to improve purity. Li2S (1.4358 g, 31.2 mmol) and Al2S3 (1.5642 g, 10.4 mmol) were weighed according to the stoichiometric 3:1 ratio. The resulting powders were then mixed and ground in an agate mortar for 15 min, transferred in an alumina crucible, and placed in a quartz tube before sealing under vacuum (10–4 mbar). The tube containing the sample was heated at 700 °C for 12 h and then to 800 °C for 12 h with an intermediate grinding step in between both firings. The heating ramp rate of the furnace was 5 °C·min–1, and cooling was performed by quenching the tube in water. Neutron powder diffraction (NPD) experiments were conducted on 7Li-enriched samples of 7Li3AlS3, using 7Li2S as the precursor material, which was synthesized according to a method described by Leube et al.[46] starting from 7Li2CO3. For consistency of the structural analysis, the 7Li3AlS3 sample was also used for synchrotron diffraction experiments.

The importance of the quenching step was investigated by either letting the sample cool down by turning off the furnace or by setting the cooling ramp rate to 1 °C·min–1. The sample cooled down by turning the furnace off did not show any extra impurity peaks in the X-ray diffraction pattern; however, the sample cooled with the 1 °C·min–1 ramp resulted in the partial decomposition into Li5AlS4 and LiAlS2. This decomposition was also observed when a formerly prepared Li3AlS3 sample was heated slowly to 700 °C at 1 °C·min-1, maintained at this temperature for 30 min, and quenched to room temperature. We therefore maintained the 5 °C·min–1 fast heating ramp rate and the quenching step for the synthesis of the material.

Probe Structure Generation and Energy Calculations

All energies were computed using periodic DFT with the VASP program.[47] The PBE functional was used[48] with the projector augmented wave approach to treat core electrons.[49]

A probe structure approach was used to sample compositions in the Li+–Al3+–O2––S2– phase space.[36] Crystal structure prediction (CSP) was used to generate a probe structure at each composition. Each probe structure was assumed to have an energy close enough to the global minimum energy structure to assess the thermodynamic stability of a hypothetical compound at that composition against the formation of an assemblage of known phases.

The CSP was performed using the in-house code ChemDASH (Chemically Directed Atom Swap Hopping). Cells containing hexagonal close-packed (hcp) and cubic close-packed (ccp) anion lattices hosting O2– and S2– were constructed, and some octahedral and tetrahedral interstitial sites occupied by Li+ and Al3+ cations. The hcp cells contained eight anions, and the ccp cell contained nine anions, with a correct number of cations to satisfy charge neutrality at each composition. Structures were initialized with a random decoration of the anion and cation sublattices, and their structures were optimized. To generate new structures, the positions of some anions were swapped on the anion sublattice, or in the interstitial sites, the positions of some cations were either swapped or moved to previously vacant interstitial sites. The new structure was then optimized by relaxation of the atomic positions to the nearest local minimum. At each step, a Monte Carlo sampling algorithm was used to accept swaps, which lowered the energy of the system or increased it by an amount lower than the Monte Carlo energy threshold. The process was continued until 1000 structures had been generated and the lowest energy structure at each composition was taken forward for calculating the stability.

One of the features of ChemDASH is to perform structural optimizations in a number of stages, which can use different parameters. This was done when optimizing each of the structures generated during the CSP process, with each stage using an increasing level of accuracy. In the first stage of each geometry optimization, Γ point-only calculations were used with a plane wave cutoff of 400 eV. By the final stage of each geometry optimization, a 2 × 2 × 2 k-point grid was used with a plane wave cutoff of 600 eV. The cell vectors and atomic positions were optimized until forces fell below 0.02 eV·Å–1.

Once a probe structure had been obtained, its energy was recalculated at a more accurate level, which was also the level of accuracy used to calculate the energies of previously reported phases. These energies were used to generate the convex hulls of chemical stability. A plane wave cutoff of 700 eV was used with a k-point spacing of 0.15 Å–1. Cell vectors and atomic positions were optimized until forces fell below 0.001 eV·Å–1. The convex hull of chemically stable compositions was generated using pymatgen.[50]

DFT calculations were also performed on an ordered analogue of the experimentally refined crystal structure of Li3AlS3. The split Li sites were merged onto a single high-symmetry site, and all sites were given full occupancy. The structure was then optimized in VASP using the more accurate parameters detailed above. No imaginary frequency modes were found in phonon calculations, showing that the structure is stable against displacement of ions from their relaxed positions. The computed phonon frequencies are presented in Table S2 of the SI.

Elemental Analysis

Elemental analysis of Li3AlS3 was performed by Mikroanalytishes Labor Pascher at Remagen-Bandorf, Germany, after dissolution in a HF/HCl solution at elevated temperature and pressure.

Diffraction

X-ray Diffraction

Synchrotron X-ray diffraction (SXRD) was performed at Diamond Light Source UK, on high-resolution beamline I11, at λ = 0.82465 Å. The sample was introduced into a 0.7 mm diameter borosilicate glass capillary to record the pattern in transmission mode [0 ° < 2θ < 150°, and Δ(2θ) = 0.004°] using a high-resolution multianalyzer crystal (MAC) detector. The experiment was performed at room temperature. For the Rietveld refinement, the Thompson–Cox–Hastings function[54] with spherical harmonic expansion implemented in FullProf was used to model the peak shape anisotropy.[55] In particular, the (−311), (−402), (−602), and (331) reflections showed a pronounced anisotropic peak shape broadening, which could be due to the presence of structural defects such as antiphase domains and stacking faults.[56] The improvement of the fit thanks to the use of spherical harmonics is illustrated in Figure S3.

Neutron Diffraction

Time-of-flight neutron powder diffraction (NPD) data were collected on Li3AlS3 using a high-resolution powder diffractometer (HRPD) instrument at ISIS, UK. Experiments were carried out at ambient temperature on the 7Li-enriched sample sealed in thin-walled vanadium cans with a diameter of 8 mm, sealed with an indium gasket under 1 atm of helium gas. For the Rietveld refinement, all banks were fitted simultaneously with the TOF pseudo-Voigt back-to-back exponential function with spherical harmonic expansion as for the SXRD data refinement. The structure determination was performed using Jana2006[68] in order to utilize the charge flipping method implemented in the software to yield an initial model. FullProf[70] then had our preference for complete refinement.

Nuclear Magnetic Resonance (NMR) Spectroscopy

The 6Li magic-angle Spinning (MAS) NMR spectra were recorded at 9.4 T on a Bruker DSX spectrometer using a 4 mm HXY MAS probe (in double resonance mode) and at 20 T on a Bruker NEO spectrometer using a 3.2 mm HXY MAS probe (in triple resonance mode). The 6Li MAS spectra were obtained at 9.4 T with a pulse length of 3 μs at a radiofrequency (f) field amplitude of ω1/2π = 83 kHz and a MAS rate of ωr/2π = 10 kHz and at 20 T with a pulse length of 4.5 μs at an rf field amplitude of ω1/2π = 56 kHz and a MAS rate of ωr/2π = 20 kHz. The 27Al MAS NMR data were recorded at 9.4 T on a Bruker Avance III HD under MAS at a rate of ωr/2π = 12 kHz using a 4 mm HXY MAS probe (in double resonance mode) and at 20 T on a Bruker NEO spectrometer using a 3.2 mm HXY MAS probe (in triple resonance mode). The 27Al spectra were obtained at 9.4 T with a short pulse angle of 30° of 0.33 μs duration at an rf amplitude of ω1/2π = 83 kHz and at 20 T with a short pulse angle of 30° of 0.55 μs duration at an rf amplitude of ω1/2π = 50 kHz. The 27Al triple quantum magic-angle spinning (MQMAS)[57] was obtained at 9.4 T with a z-filtered sequence[58] and using rf field amplitudes of ω1/2π = 83 kHz for the excitation and reconversion pulses and 4 kHz for the selective 90° pulse. All spectra were collected at room temperature and obtained under quantitative recycle delays of more than 5 times longer than the spin–lattice relaxation times T1, which were measured using the saturation recovery pulse sequence and fitted with a stretch exponential function of the form 1 – exp[−(τ/T1)β] (with β ranging from 0.3 to 1). The 6Li and 27Al shifts were referenced to 10 M LiCl in D2O and 0.1 M Al(NO3)3 in H2O at 0 ppm, respectively.

AC Impedance Spectroscopy

A pellet of the Li3AlS3 powder was made by uniaxially pressing ∼30 mg of powder in a 5 mm diameter cylindrical steel die at a pressure of 125 MPa, followed by sintering in an evacuated quartz tube at 800 °C for 12 h. A relative density of 80% was obtained by this method.

AC impedance measurements were performed using an impedance analyzer (Solartron 1296 dielectric interface coupled with a Solartron 1255B frequency response analyzer) in the frequency range from 1 MHz to 100 mHz (with an amplitude of 50 mV). Silver paint (RS silver conducting paint 186-3600), brushed on both sides of the pellet and dried under vacuum at room temperature, was used as ion blocking electrodes. Variable temperature conductivity measurements were carried out under argon (flow rate 50 mL·min–1), using a custom-built sample holder, in the temperature range 25–125 °C. The impedance spectra were fitted with an equivalent circuit using the ZView2 program.[59]

Results and Discussion

Computational/Experimental Study of the Li–Al–O–S Phase Field

We explored the Li–Al–O–S phase field using probe structure-based material discovery.[35,36] This method involves identifying a set of unexplored compositions on the phase field of interest and computationally generating a probe structure for each one. By determining the energy of each of the probe structures compared to the convex hull, we can identify low-energy regions of the phase diagram and hence target synthetic efforts toward regions of the phase field where new compounds are more likely to be found. To ensure fully occupied anion sublattices, the compositions of unit cells used in calculations were constrained to contain eight or nine anions with a stoichiometric number of cations to satisfy charge neutrality. Under these constraints, a range of compositions were chosen to span the phase field, with more compositions at the Li rich end (Figure a). We sampled all of the possible compositions for the cells containing eight anions and then only considered the Li/Al ratios closest to 1:1 for the cells containing nine anions since these were the lowest energy regions following the initial screen. All possible S/O ratios were considered. The computed energy of the probe structure at each computed composition is presented in Figure a.

Figure 1

(a) Calculated energy of different compositions in the Li–Al–O–S phase field using cells containing hexagonal close-packed (hcp, black triangles) and cubic close-packed (ccp, black filled circles) anion lattices. Ehull is the energy above the convex hull. Reported oxide and sulfide phases in the Li–Al–O–S phase field (black rectangles). (b) First-stage experimentally tested compositions, which resulted in a mixture of already reported compounds (empty red squares with black letters), and a mixture of already reported compounds along with the presence of the new phase (filled red squares with white letters). Second-stage experimentally tested compositions (numbered black circles). Composition of points are as follows: A (Li3Al9O2S13), B (LiAlOS), C (LiAlO0.2S1.8), D (LiAlO1.8S0.2), E (Li7Al2O4S), F (Li5AlO3S), 1 (Li4Al2O2S3), 2 (Li6Al8O10S5), 3 (Li2Al4O4S3), 4 (Li2Al4O5S2), and 5 (Li3AlS3).

A valley of lower energy is observed at a Li/(Li + Al) ratio of 0.5 (LiAlS2–LiAlO2 solid solution line), with a local energy minimum at 42 meV·atom–1 above the convex hull for the previously proposed composition of LiAlOS.[37] When using the full convex hull construction, LiAlOS is predicted to decompose into LiAl5O8, LiAlS2, and Li2S. In comparison, the probe structure in the previous study was also determined to be thermodynamically unstable, but with respect to LiAlS2 and LiAlO2, with a calculated formation reaction energy of 46 meV·atom–1 at 0 K.[37] Because these calculated energies are relatively low, LiAlOS could be a metastable structure that is possibly synthesizable and was therefore selected as a candidate composition for experiments.

Near the sulfur- and oxygen-rich regions of the LiAlS2–LiAlO2 solid solution line, the calculated energy is interestingly low compared to other regions close to the borders in the overall phase diagram. For instance, the energy of composition LiAlS1.8O0.2 is 34 meV·atom–1 above the convex hull. This value is close to the value of the known phase LiAl5S8 (38 meV·atom–1 above the convex hull), which has been synthesized in the literature.[60] The composition LiAlS1.8O0.2 was therefore selected for experimental synthesis. The symmetrical composition near the oxide end member, LiAlS0.2O1.8, shows a higher energy (51 meV·atom–1) but is close to that of LiAlOS and also synthesized as a matter of comparison.

Plateaus with energies below 60 meV·atom–1 above the convex hull are observed in the regions close to the terminal Al2S3 and Li2O. Within the former, the composition Li3Al9O2S13 was previously suggested by the Materials Project[61] as an interesting analogue candidate of Ga9Tl3O2S13.[62] The energy calculated for this compound in this work is 46 meV·atom–1 above the convex hull, which is also close to that calculated for LiAlOS. The plateau close to Li2O, as well as presenting relatively low energies, attracted our interest due to the high content of lithium, possibly favorable for attaining higher conductivities. Two compositions, Li5AlO3S and Li7Al2O4S, were selected for experimental synthesis in this region of the phase diagram.

Figure b and Table S1 summarize the six compositions (labeled A to F and represented by red squares) that were chosen, in a first stage, after analysis of the results from calculations. The experimental procedures for all samples were the same and were described in the Experimental Section. In particular, we used Li2O and LiAlO2 as oxide precursors whenever the composition enabled them, in order to improve the reactivity compared to the use of Al2O3. Moreover, a relatively high temperature of 800 °C and cooling by a quenching procedure were chosen in order to both facilitate the access to high-energy phases and to kinetically trap them and prevent the decomposition into binary or ternary phases during cooling.

Compositions highlighted with the empty red squares and black letters in Figure b consisted of a mixture of only already reported phases after annealing, which is detailed in Table S1. However, two points along the LiAlO2–LiAlS2 solid solution line (Figure b, points B and C) showed the presence of an unknown phase along with the already reported phases. Figure S1 shows the XRD patterns of samples A to F after reaction made in the first stage. For composition LiAlOS, the unknown phase seemed to be in a relatively high amount; thus, in a second stage, four new compositions were tested around this point (black circles in Figure b, composition given in the caption, and Table S1). For composition number 4 in particular (Li2Al4O5S2), this new phase was present along with Al2O3 as a single impurity. Figure S2 shows the XRD patterns of samples 1 to 4 after the reaction made in the second stage. Li2Al4O5S2 can then be written as Li2Al4O5S2 = a Al2O3 + Li2Al2/3(1+OS2 (0 ≤ y < 5), highlighting the fact that the unknown phase must have a Li/S ratio of 1. By considering the two end members (y = 0 and y = 5), this result led us to investigate the solid solution line x Li2Al4O5S2 + (1 – x) Li2Al2/3S2, on which the composition Li2Al2/3(1+OS2 is located. The synthesis of the pure sulfide end member (point number 5) gave the phase pure new compound sought, which was thus revealed to be a pure sulfide material, rewritten as Li3AlS3. The energy calculated for Li3AlS3 was found to be 16 meV·atom–1 above the convex hull. No other calculated energies in the phase diagram, which do not already correspond to known phases, drop below this threshold. We therefore concluded that no other oxysulfide phases are likely to form in the Li–Al–O–S field. The oxide analogue Li3AlO3 could not be stabilized by a similar quenching method (cf. Supporting Information) and leads to a calculated energy of 27 meV·atom–1 above the convex hull. The calculated energy of Li3AlS3 is slightly above the hull, which suggests that the phase is metastable.

In the Li–Al–S phase field, phases with compositions LiAlS2,[63] Li5AlS4,[64,65] and LiAl5S8[60] have previously been reported. Moreover, an amorphous phase with composition Li3AlS3 was identified previously as the discharge product of an Al–S battery.[66] Although no experimental structural data was presented for this model, a local structure of this amorphous phase was modeled with DFT and displayed a network of isolated 4- and 5-coordinated aluminum. Both the LiAlS2 and Li5AlS4 structures possess nearly close-packed anion layers arranged in a hexagonal stacking sequence. LiAl5S8 is dimorphic and comprises a low-temperature modification with a normal spinel-type structure and a high-temperature modification related to the ZnIn2S4 structure. Both LiAl5S8 structures have a cubic close-packed arrangement of the anion lattice. The integrated computation–experiment approach described here enabled the identification of a new crystalline phase with the composition Li3AlS3. This phase was isolated by synthesis, and its structure and lithium transport properties were experimentally investigated.

Synthesis and Structure of Li3AlS3

Synthesis

Polycrystalline Li3AlS3 was synthesized via a solid-state reaction of Li2S and Al2S3 (described in the Experimental Section). The powder XRD profile of the product could be indexed to a phase whose structure does not match any of the compounds previously reported for Li–Al–S systems. A small quantity of Li5AlS4 (3.3(5) mol %) was also identified in the XRD pattern. Elemental analysis gave an overall composition of Li3.1(1)Al1.1(1)S3.0(1) (Table S3), but the ICP measurement is not sufficiently precise to distinguish between the nominal reaction stoichiometry (Li3AlS3), the measured composition, and the presence of the secondary Li5AlS4 phase, further complicating its interpretation; consequently, the compound is referred to as Li3AlS3 hereafter.

Structure Determination

The crystal structure of Li3AlS3 was solved by first indexing the SXRD pattern using the first 22 reflections using GSAS-II.[67] The unit cell was indexed in the space group C12/c1 with approximate lattice parameters of 14.3 × 12.0 × 6.6 Å with β ≈ 117°. The lattice parameters were then refined across the d spacing range 16–0.67 Å (2θ = [3–75°]) via a Le Bail fit in Jana2006.[68] The structure was solved initially by locating the S and Al atoms using superflip[69] implemented in Jana2006. From this solution, a Rietveld model was refined against the SXRD and NPD patterns. Once converged, the Li atoms were located using Fourier Difference mapping on the NPD patterns, searching for peaks in the difference map located at greater than 1 Å from existing atoms within the model. When no more new sites could be located, an initial Rietveld model was constructed and refined.

Final Rietveld refinement of the neutron and synchrotron diffraction data was carried out using the program FullProf.[70] First, the SXRD pattern was fitted on its own. Position, site occupancy factor (sof), and atomic displacement parameters (adp) of lithium atoms remained fixed in the refinement of the synchrotron data. The sof of aluminum refined to 0.974(1) on its site and significant drops in the conventional reliability factors were obtained: RBragg decreased from 2.79 to 2.68. Figure a shows the final fit of the SXRD pattern, and the following results were obtained for the cell parameters: a = 14.31901(5) Å, b = 11.98037(3) Å, c = 6.62700(2) Å, and β = 116.9231(3)°. The values of the refined cell parameters were then implemented and fixed in the refinement of neutron data. The final model of the structure was obtained through the combined refinement of the neutron data coming from the three neutron datasets. Figure b–d shows the final fits of the patterns. Refinement details and outcomes as well as the crystallographic data are summarized in Tables S4 and S5.

Figure 2

Final Rietveld refinement of (a) the synchrotron X-ray diffraction pattern of 7Li3AlS3 (Diamond Light Source, I11 beam line) with fixed Li positions and (b) 7Li3AlS3 against neutron powder diffraction data (ISIS neutron source, HRPD) from (b) bank 1 (2θ = 168.330°), (c) bank 2 (2θ = 89.580°), and (d) bank 3 (2θ = 30.000°), with Iobs (red dots), Icalc (black line), Iobs – Icalc (blue line), and Bragg reflections (red tick marks for Li3AlS3, black tick marks for Li5AlS4, and blue tick marks for the vanadium can).

The unit cell shows three sulfur sites (S1, S2, and S3), three tetrahedral sites on general positions 8f (Table S5), two of them occupied by Li (Li1 and Li4) and one by Al and Li ions (Al and LiAl), and two occupied octahedral lithium sites (Li2 and Li3) located on two 4e Wyckoff positions on the 2-fold rotation axis (at the beginning of the refinement). The isotropic adp (Biso) were refined to large values: 3.3(5) Å2 for Li2 and 3.7(5) Å2 for Li3. In order to improve the model, the displacement was modeled as anisotropic, which led to the decrease of χ from 1.74, 6.46, and 1.54 to 1.73, 6.06, and 1.54 for banks 1, 2, and 3, respectively. The anisotropic ADPs for Li2 and Li3 remained high, and a marked anisotropic displacement along the a axis was found for Li3 in particular, as highlighted by the Fourier difference map in the ab plane in Figure S4a, thereby prompting us to consider site splitting. Li3 was moved from its fully occupied 4e position (0, y, 0.25) to a half occupied general 8f position (x, y, z), which then generates a second Li3 atom of coordinate (−x, y, z) from the other side of the rotation axis, within the same coordination polyhedra. For Li2, as the anisotropic displacement was not as straightforwardly along one single direction, moving the atom to one half occupied 8f position did not improve the fit. This site was split into two 8f positions (Li2 and Li2b), each allocated first with an occupancy of 0.25, generating 4 Li positions within the same coordination polyhedron, in order to model its large displacement. The occupancies of Li2 and Li2b were then refined by constraining their sum to be equal to 0.5. In that way, the total occupancy within each polyhedron is equal to one. The site splitting of Li2 and Li3 led to much smaller isotropic adp of 1.2(7) Å2 for Li2 (and Li2b) and 1.7(4) Å2 for Li3, along with a reduced residual density in the Fourier difference map around the sites (Figure S4b). Another indication for preferring the modeling of both sites with multiple atoms was the increase in the bond valence sum (BVS) from 0.76(1) to 0.83(16) for Li2 and 0.82(4) for Li2b, while keeping it to 0.80(3) for Li3.

Through occupancy refinement, a lithium antisite defect was identified and occupied 6.7% of the aluminum site. No other Li or Al antisite defects were found. Details of the refinement procedure are presented in the Supporting Information, page 7, and results of the final refinement of the occupancies are given in Table S4.

NMR spectroscopy at various fields was deployed to further confirm the overall pattern of site occupancy of the lithium atoms. The 6Li MAS NMR spectra at 9.4 and 20 T for Li3AlS3 are shown in Figure a and display three well-resolved resonances at 1.7, 1.3, and −0.2 ppm, which fit yield signals of equal integration. A small shoulder is also observed at 1 ppm (Figure S8) and corresponds to the Li5AlS4 impurity seen in the diffraction and 27Al NMR data (see below);[65] this signal was found to integrate 3.0(5) mol % Li3AlS3, in agreement with the 3.3(5) mol % value from diffraction. Based on the well-established semi-empirical correlations relating the lithium coordination environment and 6Li NMR shifts,[71] further aided by calculations of the NMR parameters using the GIPAW approach[52,53] as implemented in CASTEP[51] (cf. Supporting Information, Table S7), the signal at −0.2 ppm has been attributed to the octahedrally coordinated Li2/Li2b and Li3 sites while the resonances at 1.3 and 1.7 ppm correspond to Li4 and Li1, respectively. These assignments agree well with the structural refinement described above, which identified the sum of the contents of the three octahedrally coordinated sites Li2 (0.8(3)), Li2b (3.2(3)), and Li3 (4.0(4)) to 8.0(7) Li per unit cell and the contents of the two tetrahedrally coordinated Li4 and Li1 to 7.8(3) and 8.00 per unit cell, respectively.

Figure 3

(a) 6Li MAS spectrum of Li3AlS3 obtained at magnetic fields of 9.4 T (black) and 20 T (blue). The experimental spectrum (full lines), total fit (dashed lines) spectral deconvolution (dotted lines), Li5AlS4 impurity (red dotted lines), and GIPAW-simulated spectrum (green lines) are shown. (b) 27Al MQMAS NMR spectrum of Li3AlS3 recorded at a magnetic field of 9.4 T and 20 T. The dotted lines (black for a field of 9.4 T and blue for 20 T) and the red dotted lines represent the spectral deconvolution of Li3AlS3 and Li5AlS4, respectively. The dashed lines show the total fit for the sample, and the solid lines show the anisotropic one-dimensional 27Al spectrum, while the vertical spectrum shows the non-quantitative isotropic 27Al spectrum. The solid green line shows the GIPAW-simulated spectrum with an isotropic chemical shift of 117 ppm, a quadrupolar coupling constant of 5.1 MHz and an asymmetry parameter of 0.44 (Table S7).

The 27Al one-dimensional MAS spectra for Li3AlS3 at 9.4 and 20 T are shown in Figure b and reveal a second order quadrupolar line shape that resonates at ∼100 ppm (at 9.4 T) and ∼120 ppm (at 20 T), typical of tetrahedrally coordinated Al sites, and a very sharp signal, which is field-independent, at ∼130 ppm, corresponding to the small amount of Li5AlS4 impurity (Figure S7). Note that no signal in the octahedral region (around 0 ppm) of the 27Al MAS NMR spectrum is present as expected. The z-filtered triple quantum MAS[53,54] NMR spectrum of Li3AlS3 is also shown in Figure b and demonstrates that the ∼100 ppm signal corresponds to one resonance only with an isotropic chemical shift of 117 ppm (cf. the Supporting Information) and a quadrupolar coupling constant (CQ) of 5.8 MHz in close agreement with the computed value of 5.1 MHz (cf. SI, Table S7).

The final model led to the overall refined composition Li3.13(2)Al0.958(4)S3, slightly different from the composition determined by ICP. The small differences can be explained by the presence of the Li5AlS4 impurity, which prevents the ICP measurement from producing an accuracy to the nearest two decimal places. The refined composition is different from the ideal, which will be discussed hereafter.

Structure Description

Polyhedral Arrangement

Li3AlS3 adopts a structure related to that of Na3InS3 reported by Eisenmann and Hofmann where In3+ and Na+ cations are replaced by Al3+ and Li+, respectively.[72] The only other isostructural phases reported are Na2Mn2S3,[73] Na2Mn2Se3,[74] and LiNa1–Mn2S3 (x ≈ 0.7)[75] in which half of the Mn2+ ions are replaced by aluminum atoms whereas the other half along with the sodium atoms are replaced by Li+ ions. The structure is constructed from an hcp arrangement of sulfur atoms with an AB A*B* stacking of anion layers where B is the equivalent of A through the c glide plane and 2-fold axis symmetry operations. A* and B* are the equivalent of A and B through the C centering translation (Figure a). In the tetrahedral layer, Li1 and Al atoms occupy 2/3 of the tetrahedral interstices between a pair of sulfur atom layers (B and A*, equivalent to the B*A pair). Between the second pair of sulfur layers (A and B, equivalent to the A*B* pair), Li2 and Li3 occupy octahedral interstices, whereas Li4 occupies a tetrahedral interstice, generating a mixed polyhedral (octahedral–tetrahedral) layer. The two different polyhedral layers are stacked alternately perpendicular to the bc plane (Figure a). Bond distances and angles of the different polyhedral and BVS calculations performed for all atoms are summarized in Table S6.

Figure 4

(a) Crystal structure of Li3AlS3 showing the alternating arrangement perpendicular to the bc plane of the tetrahedral layers containing AlS4 and LiS4 tetrahedra and the mixed polyhedral layers containing Li-only polyhedra. (b) T+ and T– interstices in the tetrahedral layer, showing the corner-sharing arrangement of the Li1, Al, and vacant (empty) tetrahedra in each network, as well as the interconnection (following the yellow arrow) of each T+ (thin lines) and T– (thick lines) network so that AlS4 dimers are formed. The highlighted yellow face of the Li1 tetrahedron corresponds to the only face that shares two edges with two vacant sites. (c) View of both the mixed polyhedral layer and the tetrahedral layer in the bc plane and of their interconnection (following the yellow arrow). Polyhedra colors: blue: Al tetrahedra; orange: Li tetrahedra; red: Li2 octahedra; light red: Li3 octahedra.

In the tetrahedral layer, one tetrahedral site in every three is vacant in an ordered manner (Figure b). Each T+ and T– interstice forms a network of alternating Al, Li1, and vacancy corner-shared tetrahedra. The T+ and T– networks interlock in such a way that each AlS4 unit shares one edge (S3–S3) with another Al tetrahedron and a second edge with the Li1S4 unit within the layer. These 4-edge-shared tetrahedra (Li1S4(T+)–AlS4(T–)–AlS4(T+)–Li1S4(T+)) form a unit that is connected to other units of this type by corner-sharing (Figure a, circled part in Figure b). The Li1 tetrahedron has two S atoms that share corners with Li1 and Al, and there is no corner-sharing of Al tetrahedra.

Figure 5

Coordination polyhedra of (a) Li1 and Al in the tetrahedral layer, (b) Li4, (c) Li2 and Li2b, and (d) Li3 in the mixed polyhedral layer.

In the mixed polyhedral layer, each of the Li2, Li2b, and Li3 atoms is surrounded by six sulfur atoms to form LiS6 octahedra (Figures c and 5b). Li2 and Li2b polyhedra form an infinite chain of edge-shared octahedra along the c axis, which are also connected to three Li3 octahedra also through three shared edges. Thus, both octahedra form infinite two octahedra-wide chains along the c axis. Each chain is separated along the b axis by a chain of edge-shared LiS4 T+ and T– tetrahedra occupied by the Li4 atom (Figure c). Li3 octahedra are linked to two consecutive T+ and T– Li4 tetrahedra, sharing one face with each of them (Figure c). These two Li4 tetrahedra form a unit (Figure b) and each T+ and T– of the unit shares one corner with one Li2/Li2b octahedron from the same chain as the face-shared Li3 octahedra, as well as one corner with one Li3 octahedron of the chain on the other side of the Li4 chain. Along the Li4 chain, octahedral interstices are vacant so that only 2/3 of the octahedral sites are occupied in the layer. As a result of crystallization, the ordered structure obtained experimentally in this study strongly differs with the structure modeled for the amorphous discharged product with the same composition described by Yu et al.[66]

Figure c shows the connection between the polyhedra of both layer types. Each of the Al and Li1 T+ (T–) tetrahedra is connected to the layer below (above) by sharing the face at the base of the tetrahedron with a vacant tetrahedral site of the mixed polyhedral layer. For the Al tetrahedra, this face shares one edge with the Li2 octahedra and another edge with the Li3 octahedra of the same chain, while the third edge is shared with the vacant octahedral site. For the Li1 tetrahedra, this face shares two edges with the Li2 octahedra and one edge with the Li3 octahedra of the same chain. The S3 atom, which is at the vertex of the Al T+ (T–) tetrahedra, is shared with one Li2, one Li3, and four Li4 of the above (below) mixed polyhedral layer, whereas the S1 atom, which is at the vertex of the Li1 T+ (T–) tetrahedra, is shared with one Li3 octahedron and four Li4 of the above (below) mixed polyhedral layer. The chain of vacant T+ (T–) sites along the c axis in the tetrahedral layer lies above (below) the chain of Li4 tetrahedra of the mixed polyhedral layer.

Polyhedral Distortions and Atom Displacements

Figure shows the bonding environment of each of the cations. In the tetrahedral-only layer, Li1 atoms are strongly off-centered toward the S1–S2–S2 face so that it both does not share any edges with the AlS4 tetrahedra in the same layer and does not belong to the octahedral layer (Figure ). This face is also the only one that shares two edges with two vacant sites of the tetrahedral layer as opposed to one for the other two faces that do not belong to the octahedral layer (highlighted yellow face in Figures b and 6). Moving away from the nearby Al cations as well as from the Li2 and Li3 atoms of the octahedral layer provides the electrostatic driving force for its displacement (Figure ). The BVS of Li1 is 0.97(2), which is very close to the ideal value of +1, considering the oxidation state of lithium.

Figure 6

Crystal structure of Li3AlS3 showing the arrangement of octahedral (red) and tetrahedral (orange) lithium and tetrahedral aluminum (blue). The direction of the displacement of atoms is symbolized by arrows: blue for Al, orange for Li1 and Li4, and yellow for S.

The tetrahedral Li4 site in the mixed layer shares a face with the vacant tetrahedral site of the tetrahedral layer. The Li4 position is strongly displaced toward the base of the tetrahedron and the adjacent sulfur layer, toward this vacant tetrahedral interstice (Figure ). Li4 thus adopts a pseudo-trigonal bipyramid environment with one short axial Li4–S1 and one long axial Li4–S2 bond. BVS for Li4 is slightly lower than that of Li1 (0.93(2)), which is consistent with its pseudo-5-coordinated environment, making it more loosely bound to the S atoms compared to the 4-coordinated Li1.

This is reflected in the NMR resonance frequency of Li4 (1.3 ppm) for which the 5-coordinate environment provides this site with an intermediate chemical shift between the lithium atoms occupying a tetrahedral site (Li1 at 1.7 ppm) and octahedral sites (Li2/Li2b and Li3 at −0.2 ppm). The NMR shift calculation also suggests that the resonances of Li3 and Li2/Li2b should be resolved as the isotropic chemical shifts differ by 0.3 ppm; however, this is not observed experimentally, even at high field, perhaps due to the larger full width at half-maximum observed for this Li3/Li2/Li2b resonance (18 Hz compared to 10 and 11 Hz for Li4 and Li1, respectively). The presence of the occupied Li4 sites distinguishes both Li3 and Li2/Li2b sites: Li3 shares faces with Li4, whereas Li2 does not (Figure c). The shift of the Li4 toward the adjacent sulfur layer pushes the Li4 atom further away from the octahedral Li3 atoms and therefore reduces the structural difference between Li2 and Li3, further explaining the similar shifts observed experimentally for these octahedral sites (Table S7). On the contrary, the difference in the environment of Li1 and Li4 is more pronounced (tetrahedral vs trigonal bipyramid), which explains the resolution of their respective NMR peaks. The consistency between the experimental NMR data and the computed ones from the described structure further reinforces the accuracy of the selected structural model.

The geometry of the AlS4 dimer is represented in Figure a, and among the six edges of each AlS4 tetrahedron, four of them are directly connected to the S3–S3 edge-shared with the other tetrahedron of the dimer. Consequently, the aluminum position is shifted toward the S1–S2 edge that does not share a common sulfur atom with the other edge-shared Al tetrahedron of the dimer. This displacement is symbolized by blue arrows in Figure . This can be explained by the proximity to the other Al atom of the edge-shared dimers with which it tends to maximize its distance. The Al–Al distance is 3.015(6) Å, which is similar to distances obtained in other sulfide materials (Na6Al2S6, Na3FeS3) presenting these dimers.[76,77] This is also supported by the 27Al NMR data obtained experimentally and computed with GIPAW (cf. SITable S7) that yield a clear second-order quadrupolar line shape at both 9.4 and 20 T and from which large quadrupolar coupling constants (CQ = 5.8 and 5.1 MHz for experimental and computed values, respectively) and distorted asymmetry parameter values (ηQ = 0.56 vs 0.44 for experimental and computed values, respectively) are obtained. These data clearly demonstrate the non-symmetrical dimeric Al coordination environment in Li3AlS3 and is in sharp contrast to the symmetric AlS4 tetrahedra observed in Li5AlS4 (CQ ≈ 0 MHz, Figure S7) as evidenced by the field-independent NMR narrow line of this phase.

The dimerization leads to the strong repulsive force between both highly charged Al3+ cations. This in turn brings the S3 atoms inside the dimer toward each other and therefore toward the interior of the tetrahedral layer, in order to keep the BVS of Al and S close to their ideal values (symbolized by yellow arrows in Figure b,c). The S3–S3 distance (represented by a thick line in Figure b,c, dS3–S3 = 3.463(14) Å) is indeed the shortest in the structure. The compression of the S3–S3 edge of the Al2S6 dimer unit in the tetrahedral layer leads to the stretching of the S3–S3 edge of Li2(Li2b)S6 and Li3S6 octahedra in the mixed polyhedral layer (Figure b,c).

As shown in Figure a, the Li2b (and Li2) octahedron is highly distorted, and the S3–S3 edge is considerably longer than the other edges (dS3–S3 = 4.40(1) Å whereas the length of the other edges ranges from 3.66(1) to 3.951(7) Å). Also, the Li3 octahedron is elongated along an axis defined by two S3 atoms (dS3–S3 = 5.729(14) Å), whereas the equatorial plane defined by two S1 and two S2 atoms is close to a square (dS2–S2 = 3.893(13) Å, dS1–S1 = 3.713(14) Å, dS2–S1 = 3.682(8) Å). The explanation for the site splitting of Li2 can be linked to the distribution of the sulfur vacancies (Supporting Information, page 10). In contrast to the tetrahedral Li1 and Li4, the BVS values for octahedral Li2, Li2b, and Li3 are below the theoretical value (0.83(16), 0.82(4), and 0.80(3) for Li2, Li2b and Li3, respectively), which reflects the fact that they are weakly bonded to the sulfur atoms. This observation has also been made in similar structures.[65]

A sulfur deficiency was found on two S sites. This deficiency accounts for the charge compensation with the Al defect. The majority of sulfur vacancies are located on the S3 site, which bridges the two aluminum tetrahedra of the dimer. As explained above, the S3–S3 distance is the shortest in the structure; hence, vacancies here would reduce anion–anion repulsions. Again, the presence of the dimers leading to the short S3–S3 distance could be the trigger for the deficiency of the S site. Further, the occupancy of the Al site cation vacancies by Li creates a negatively charged antisite defect, which could also drive the localization of the positively charged sulfur vacancy on the S3 site.

The analysis of the Li3AlS3 structure highlights the importance that the Al2S6 dimers have on site geometries as well as on site occupancies of both lithium and sulfur atoms. Moreover, the comparison with the probe structure, which also shows the presence of these dimers (Supporting Information, page 11, and Figure S5), further suggests that the stability of the structure is connected to the presence of the Al2S6 dimers.

Comparison with Known Structures

Lithium sulfide materials showing different arrangements of cation polyhedra in hcp arrays are common, and some interesting compositions and structures have been described by Lim et al.[65] Among them, one can cite Li2FeS2, which consists of an octahedral-only layer whose interstices are 100% occupied by lithium atoms, alternating with a tetrahedral layer, where all the tetrahedral sites are occupied by Li and Fe atoms randomly distributed in a 50:50 ratio.[78] Li5AlS4 shows a similar structure where only half of the Fe atoms are replaced by Al3+ ions whereas the other half is occupied by Li+ ions.[65] The tetrahedral layer consists of ordered LiS4 and AlS4 units in a 3:1 arrangement. The authors note that these structures can be expressed as [LiFe]T[Li]OS2 and [Li1.5Al0.5]T[Li]OS2, respectively. We added the superscripts “T” and “O”, which refer to the tetrahedral and octahedral coordination of the cation in alternating layers. Following this representation, the material reported in this study, Li3AlS3, could be written as [Li2/3OLi2/3T][Li2/3Al2/3]TS2 where the cation in the same square bracket belong to the same layer. Recently, Leube et al. reported a family of compounds Li4.4M0.4M′0.6S4 (M = Al3+, Ga3+, M′ = Ge4+, Sn4+) whose structure is closely related to that of Li5AlS4.[46] The highly charged cations share the same site in the tetrahedral layer, the octahedral layer is made of ordered partially occupied lithium sites and fully vacant sites in a 3:1 arrangement, and the remaining lithium atoms share two crystallographically distinct tetrahedral sites in both the tetrahedral and octahedral layer in a 74:26 ratio. Following the same convention, this structure can be written as [Li0.66OLi0.38T][Li1.11M0.2M′0.3]TS2.

Other alkali metal sulfides with composition A3MS3, where A is an alkali monovalent cation and M is a trivalent cation, have been reported in the literature. As stated above, Na3InS3 in particular shows a closely related structure to Li3AlS3 and can be written as [Na2/3□1/3]O[Na]2/3T[Na2/3In2/3□2/3]T′S2 where In and Na occupy the same sites as Al and Li in Li3AlS3. Moreover, both compounds show identical distortions of the alkali octahedra and the same direction of displacements for the tetrahedral alkali cations. This particularity is most probably coming from the high degree of constraint imposed by the presence of the M3+S4 dimers, and the fact that this observation is made in both phases reinforces the validity of this explanation. Rothenberger et al.[79] reported a structure with the composition (M(AlS2)(GeS2)4 (M = Na, Ag, Cu) where 20% of the Al atoms are in similar dimeric tetrahedral units). However, these Al dimers were much less distorted with very similar Al–S interatomic distances within the tetrahedra (the average Al–S distance is 2.2117(5) Å and the standard deviation 1.3%). The absence of tetrahedral distortion in these dimers can be explained by the 3D polyanionic rigid structure imposed by the three other Al or Ge sites. This lessening of the distortion from tetrahedral symmetry is also reflected in the 27Al NMR where the resonance observed  shows,on the spectra, smaller values of CQ for (M(AlS2)(GeS2)4 (M = Na, Ag, Cu).

The splitting of the Li2 site in Li3AlS3 is not reported for Na3InS3, nor is the displacement from the Wyckoff position to a general position of the second octahedral alkali site. Lithium is smaller than sodium, and the volume occupied by each atom considering a hard sphere model is 12 and 18% of the total volume of each octahedron for Li and Na, respectively. The displacements and splitting modeled for lithium might therefore be key to generating an appropriate bonding environment as defined by the BVS.

Interestingly, Na3AlS3 along with Na3GaS3 shows a slightly different structure (Figure a,b).[76,80] In between two sulfide layers, Al2 occupies 1/3 of the tetrahedral sites and forms dimers between T+ and T– tetrahedra of the same slab. In this layer, Na4 occupies 1/3 of the distorted octahedral sites forming infinite chains along the c axis. In between the next two sulfide layers, Na1 occupies one distorted octahedral site, whereas Na2 and Na3 form 5-coordinated sulfide polyhedra in a trigonal bipyramid configuration and all three Na sites are in a 1:1:1 arrangement. The next slab consists of a similar AlS4 and NaS6 polyhedra arrangement to the first described slab, with a slight tilt of the polyhedra (Figure b). Na5 lies within each of the sulfide layers forming the Al2S4 slab and is 6-coordinated to four of the sulfur atoms of the same layer, one in the layer above and one in the layer below, therefore forming a 2D network of edge- and corner-shared octahedra. The structure difference between Na3InS3 or Li3AlS3 (for which Na or Li is 4-coordinated in the tetrahedral layer) and Na3AlS3 or Na3GaS3 (for which Na is 6-coordinated in the tetrahedral layer) could be attributed to the size of the M3+ cation with respect to that of the alkali cation. Indeed, the ionic radius of Al3+ (0.39 Å) and Ga3+ (0.47 Å) is considerably smaller than that of In3+ (0.62 Å),[81] so the size of tetrahedral interstices will decrease and might not be suitable to host Na+ cations, which would then prefer to occupy octahedral sites, in contrast with the smaller Li+ cation.

Figure 7

(a) Crystal structure of Na3AlS3 showing the alternating arrangement, along a, of the tetrahedral layers containing AlS4 and NaS4 tetrahedra and of the mixed polyhedral layers containing Na-only polyhedra. (b) View of the two consecutive tetrahedral layers of Na3AlS3 in the bc plane. (c) Crystal structure of Na3FeS3 showing one type of layer along b (d) View of the layer along b of Na3FeS3 showing the fully occupied octahedral sites by Na atoms and the 1/3 occupied tetrahedral interstices by Fe atoms in a dimer arrangement.

Another known structure of similar composition showing different arrangements of metal and alkali cations within the hcp array is that of Na3FeS3,[77] which is also adopted by Na3FeSe3,[82] Na3AlSe3,[83] and Na3GaSe3[84] (Figure c). In those phases, the M3+ cation assembles in tetrahedral dimers in between each of the sulfide layers, and the sodium atom occupies all of the octahedral sites in these layers (Figure c,d). Curiously, the absence of alkali-only layers and the stabilization of a single layer type containing both Fe or M (M = Al3+, Ga3+, and In3+) and Na or Li polyhedra seem to occur in all of the selenide compounds as well as both the iron sulfide and selenide phases. It seems that this could be linked to the less ionic character of the bonds in those structures. Indeed, the higher polarizability of Se2– compared to S2– anions generally leads to softer and more covalent bonds in selenides.[85] It has also been shown, through magnetic property measurements, that the 3d orbitals of Fe3+ in Na3FeS3 are more extended than that of ionic Fe3+, which can be attributed to Fe–S covalency.[86] Because Al3+, Ga3+, and In3+ do not show this effect, the ionic character of the M–S (M = Al3+, Ga3+ and In3+) bond might be more pronounced in Na3MS3 than in Na3FeS3. Different layer-type structures are often found in materials where ions have different chemical properties, in particular different polarizability, which results in different ionicity of the cation–anion bond in each layer.[87,88] A more covalent character of the bonds would then attenuate the difference in polarizability of the cations and favor the single layer-type structure.

Lithium Conductivity

The ionic conductivity of Li3AlS3 was assessed by electrochemical impedance spectroscopy on sintered pellets with an 80 ± 2% relative density. The Nyquist plot of the sample measured at room temperature under an argon atmosphere is shown in Figure a. The presence of the two semicircles is characteristic of two unique time constants and therefore of the dissociation between different scattering contributions. The plot has therefore been fitted by a two-component equivalent electrical circuit (inset in Figure a) that models these two contributions. Each component consists of a resistance associated in parallel with a constant phase element (CPE, a modified capacitor taking into account inhomogeneities in the sample, cf. Supporting Information). The values of the capacitance obtained for the semicircles at high and low frequencies were 8(1) × 10–12 and 2.6(4) × 10–9 F and are characteristic of the bulk and grain boundary response, respectively.[89] The high-frequency intercepts of both semicircles give direct values of the bulk and total resistance (Rbulk and Rt = Rbulk + RGB, respectively, where RGB is the resistance resulting from the grain boundary scattering). Bulk and grain boundary room-temperature conductivities of 1.3(1) × 10–8 and 2.2(2) × 10–9 S·cm–1, respectively, were obtained. The impedance of the pellet was measured over the temperature range (24–125 °C), and each Nyquist plot was fitted with the described equivalent circuit. Tables S8 and S9 present the results of the fits and the values of the different parameters obtained at each temperature. For each temperature, the conductivity of the bulk, σbulk, was therefore extracted and showed to follow the Arrhenius law (Figure b) with an activation energy of 0.48(1) eV.

Figure 8

(a) Nyquist plot at 30 °C of Li3AlS3 and (inset) electrical equivalent circuit showing the two contributions to the conductivity. (b) Arrhenius plot of the bulk conductivity of Li3AlS3 measured by AC impedance. Black squares correspond to the experimental data, and the red line corresponds to the fits.

The room-temperature bulk conductivity is of the same order of magnitude as that of Li5AlS4 (σ = 9.7 × 10–9 S·cm–1).[65] In Li5AlS4, all Li sites are fully occupied, whereas in Li3AlS3, there are multiple ordered vacancy sites: one third of the octahedral interstices in the mixed polyhedral layer and two thirds of the tetrahedral interstices in the tetrahedral layer are vacant. Although the lowering of the activation energy (0.48 eV for Li3AlS3 and 0.61 eV for Li5AlS4) is indeed observed, the increase in Li mobility is not, which suggests that ordered vacancy sites are not sufficient to improve conductivity in this structure type.

In the recently reported related Li4.4M0.4M′0.6S4 compounds, which shows considerably higher conductivities σ = 10–5–10–6 S·cm–1, the presence of multiple disordered partially occupied lithium sites has been shown to play a major role in the improvement of the conductivity. In Li3AlS3, the only disordered vacancies can be found within the Li4 site, which was determined to be 98% occupied (Table S5), as all other Li sites are completely occupied. In order to highlight the different types of vacancies within the structure, Li3AlS3 can be written aswith □ and Δ being the disordered and ordered vacancies, respectively. The family of quaternary materials, using the same convention, can be written as

Overall, this structure shows fewer ordered vacant sites, but the content of disordered vacancies is considerably higher, which underlines the importance of this feature in order to improve conductivity.

It is interesting to note that, in the Na3FeS3 structure, only one type of layer in which all the sodium is located in octahedral sites is present. Because the mobile species are believed to be the octahedral lithium in these types of structures,[46] it would be of high interest to stabilize the Na3FeS3 structure in lithium-containing compounds while creating a large number of disordered vacancies.

Influence of the Al2S6 Dimers on the Structure and Li Ion Conductivity

The strong differences between Li5AlS4 or Li4.4M0.4M′0.6S4 (M = Al3+, Ga3+, M′ = Ge4+, Sn4+) structures and Li3AlS3 come from the ordering of the tetrahedral cations and the presence of edge-shared AlS4 tetrahedra pairs, which form Al2S6 dimers. The three structures have the same anion sublattice with tetrahedral Li and Al layers alternated with Li only layers (= mixed polyhedral layer). In the three structures, Al atoms must spread over the tetrahedral interstices of the tetrahedral layer, and when this is done in an ordered manner, the M/S ratio imposes the arrangement pattern. Indeed, in Li3AlS3, M/S = 1/3 means that there is one Al atom for three tetrahedral interstices (one tetrahedra contains four sulfur atoms, and each tetrahedra is connected to four other tetrahedra in a layer consisting of T+ and T– interstices, cf. Figure ). Therefore, M/S = 1/3 imposes a 1:2 ordering in the Al tetrahedra by optimizing the distance between each AlS4 unit (Figure , entry 1). In the same way, M/S = 1/4 in Li5AlS4 and Li4.4M0.4M′0.6S4 leads to a 1:3 arrangement of the M (or M′) tetrahedra in the layer (Figure S6). The interconnection of both T+ and T– networks presenting this arrangement type inevitably leads to the presence of the Al2S6 dimers in Li3AlS3 on the contrary to the other two phases.

Figure 9

Representation of the influence of the M/S = 1/3 ratio on the structure and arrangement of Li polyhedra in Li3AlS3 having the “Li5AlS4-type” structure leading to the presence of ordered vacancies in the tetrahedral layer.

These dimers have a crucial influence on the displacement of atoms within the structure. Because of the strong electronic repulsive forces between the two Al3+ cations in the dimer, the Al position is driven off the center of the tetrahedron and the S3–S3 distances are successively compressed and stretched among the layers (Figure , entry 2). In Li5AlS4 and Li4.4M0.4M′0.6S4, on the other hand, the S–S edges are not as compressed. This is reflected in the values of the standard deviations for the S–S distances: 0.13 (Li5AlS4), 0.20 (Li4.4Al0.4Ge0.6S4), and 0.24 (Li3AlS3) and in the maximum/minimum S–S distances (in Å): 4.1564/3.6805 (Li5AlS4), 4.3176/3.6016 (Li4.4Al0.4Ge.6S4), and 4.4206/3.4631 (Li3AlS3). This strongly impacts the geometry and the volume of the polyhedral sites in the adjacent layer, i.e., the Li-only layer (Figure , entry 3). In Li3AlS3, the volumes of the three different octahedral interstices are 26.7227, 27.1162, and 30.4551 Å3. The latter is considerably bigger than the other two and is not favorable to hosting small Li atoms. This explains why only 2/3 of the octahedral sites are occupied in the Li-only layer (Li2, Li2b, and Li3). In the same way, in Li4.4Al0.4Ge0.6S4, 3/4 of the octahedral sites are occupied, and the volume of the ordered vacant site is 31.5654 Å3. In Li5AlS4, all of the octahedral sites have the adequate geometry to host lithium, therefore leading to a full occupation of the octahedral sites within this layer. In Li3AlS3, it will be preferable for the remaining Li atoms to occupy a tetrahedral site. Li4 thus occupies the tetrahedral interstices, which are away from the already occupied octahedra as well as away from the above (or below) occupied Al tetrahedra (Figure , entry 3). The positions of the Li4 atoms are driven toward the above (or below) S–S slab to minimize repulsion with the other edge-shared Li4 tetrahedra, creating a pseudo-bipyramid trigonal environment. The displacement of the Li4 atoms is then detrimental to the occupation by another cation of the above or below tetrahedra in the tetrahedral layer. The latter would then more preferably be left vacant. This explains why among the two other tetrahedral sites not occupied by Al, only one is occupied by Li, and a 1:1:1 arrangement of Al, Li, and ordered vacancies is stabilized (Figure , entry 4). In Li4.4Al0.4Ge0.6S4, apart from the tetrahedral site that lies just below the MS4 tetrahedra of the above tetrahedral layer, all the other tetrahedral sites of the Li-only layer are equivalent and Li atoms therefore randomly occupy these positions. The motivation for the ordering of vacancies in the tetrahedral layer is therefore suppressed, and the remaining Li ions are delocalized among all tetrahedral sites in this layer. The ordered vacancy sites, unfavorable to host lithium atoms, are likely to act as a barrier for Li diffusion, and result in lower lithium ionic conductivity in Li3AlS3 than in Li4.4M0.4M′0.6S4. A complete study of the Li energy landscape to elucidate Li diffusion pathways will be undertaken to yield further insight into the role of these structural features.

In this structure type, the M/S ratio imposes the arrangement pattern of the Al tetrahedra in dimers or isolated tetrahedra and triggers the stabilization of ordered or disordered vacancies within both the tetrahedral and the Li-only layer. The lithium conductivity properties of each of the compounds can then directly be related to the structure through the amount of disordered Li vacancies, which itself can be explained by the composition through the M/S ratio. This work illustrates that, in this structure type, an M/S ratio that is too large causes structural arrangements that inhibit the diffusion of Li.

Conclusions

We have investigated the Li–Al–O–S phase diagram through a probe structure approach, which combined experimental and computational studies to yield a new phase while ruling out others. Indeed, this study revealed that no oxysulfide phases could be successfully obtained, but led to the discovery of the new sulfide Li3AlS3. The structure and properties of this compound were determined by means of high-resolution X-ray and neutron diffraction, multinuclear NMR spectroscopy at various fields, and electrochemical impedance spectroscopy. The stability of the new phase is believed to rely on the presence of AlS4 dimers, a peculiar feature not observed before in other Li ion conducting phases. The structure was described by comparing the cation polyhedral arrangements with those of other related phases with similar compositions, such as Na3MCh3 (M = Al, Ga, In; Ch = S, Se), Li2FeS2, and Li5AlS4. The study of this new compound in comparison with other similar sodium and lithium chalcogenide phases widens the spectra of possibilities to explore new interesting structures in related phase fields.