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LiPSI: A Li Superionic Conductor synthesized by a Solvent-based Soft Chemistry Approach Stefan J. Sedlmaier, Sylvio Indris, Christian Dietrich, Murat Yavuz, Christoph Dräger, Falk von Seggern, Heino Sommer, and Jürgen Janek Chem. Mater., Just Accepted Manuscript • DOI: 10.1021/acs.chemmater.7b00013 • Publication Date (Web): 23 Jan 2017 Downloaded from http://pubs.acs.org on January 25, 2017

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Chemistry of Materials

Li4PS4I: A Li+ Superionic Conductor synthesized by a Solvent-based Soft Chemistry Approach Stefan J. Sedlmaier,†,* Sylvio Indris,‡ Christian Dietrich,$ Murat Yavuz,‡ Christoph Dräger,‡ Falk von Seggern,#,§ Heino Sommer,€ Jürgen Janek†,$,* † BELLA Battery and Electrochemistry Laboratory, Institute of Nanotechnology (INT), Karlsruhe Institute of Technology, Hermann-von-Helmholtz-Platz 1, D-76344 Eggenstein-Leopoldshafen, Germany ‡ Institute for Applied Materials – Energy Storage Systems (IAM-ESS), Institute for Applied Materials, Karlsruhe Institute of Technology, Hermann-von-Helmholtz-Platz 1, D-76344 Eggenstein-Leopoldshafen, Germany $ Institute of Physical Chemistry, Justus-Liebig-University Giessen, Heinrich-Buff-Ring 17, D-35392 Giessen, Germany § KIT-TUD Joint Research Laboratory Nanomaterials, Institute of Materials Science, Technical University of Darmstadt (TUD), D-64287 Darmstadt, Jovanka-Bontschits-Strasse 2, Germany # Institute of Nanotechnology (INT), Karlsruhe Institute of Technology, Hermann-von-Helmholtz-Platz 1, D-76344 Eggenstein-Leopoldshafen, Germany € BASF SE, Carl-Bosch-Strasse 38, D-67056 Ludwigshafen am Rhein, Germany

ABSTRACT: A novel crystalline lithium superionic conductor, Li4PS4I, has been discovered utilizing a solvent-based synthesis approach. It was found that the starting material Li3PS4·DME reacts with LiI in a 1:1 ratio in DME to give a precursor that results in Li4PS4I after soft heat treatment at around 200°C in vacuum. Its crystal structure was solved ab initio by evaluating both, X-ray (Mo-Kα1) and neutron (TOF, GEM, ISIS) powder diffraction data, in a combined refinement (P4/nmm, Z = 2, a = 8.48284(12) Å, c = 5.93013(11) Å, wRp = 0.02973, GoF = 1.21499). The final structure model comprises, besides Li+ ions, isolated PS43− tetrahedra in a layer-like arrangement perpendicular to the c-axis that are held apart by I− ions. The Li+ ions are distributed over five partially occupied sites residing in 4-, 5- and 6-fold coordination environments. A topostructural analysis of the voids and channels within the PS4I4− substructure suggested a 3-dimensional migration pathway system for the Li+ ions in Li4PS4I. The Li+ ion mobility was studied by temperature-dependent impedance spectroscopy as well as 7Li solid-state nuclear magnetic resonance (NMR) spectroscopy including the measurement of spin-lattice relaxation rates T1–1. The total ionic conductivity was determined to be in the range of 6.4 · 10–5 to 1.2 · 10–4 S·cm–1 at room temperature with activation energies (EA) of 0.37 to 0.43 eV. The NMR analyses revealed a hopping rate of the Li+ ions of τ–1 = 5 · 108 s–1 corresponding to a bulk conductivity of 1.3 · 10–3 S·cm–1 at 500 K and an activation energy EA = 0.23(1) eV.

INTRODUCTION All-solid-state secondary batteries (ASSB) are an emerging option for next-generation lithium ion battery technologies. They have the potential of outperforming conventional lithium ion batteries in terms of safety, lifetime and, most importantly, energy density to significantly increase e.g. the driving range of electric vehicles. However, the whole potential of ASSBs can only be exploited if inorganic materials are identified that are able to adequately replace liquid electrolytes. Those solid electrolytes are required to robustly perform in ambitious battery systems with high-voltage cathode materials such as NCM (LiNi1/3Co1/3Mn1/3O2) or LNMO (LiNi1/2Mn3/2O4) and accordingly to exhibit properties like a sufficient Li+ ion conductivity (>10−4 Scm−1), chemical compatibility with cathode and anode materials, and electrochemical stability over a wide voltage range.1-4

Over decades, numerous lithium ion conducting materials have been discovered within different compound classes and structure types.5 Oxide materials, mostly appearing in the perovskite (Li3xLa2/3-2x□1/3-2xTiO3),6 NASICON (Li1+xTi2and garnet structure type (e.g. xAlx(PO4)3(x = 0; 0.3)), Li6.5La2.5BaZrTaO12),7 usually exhibit great chemical, mechanical, and electrochemical stability and show reasonable ionic conductivities up to 10−3 S cm−1, however their rigidity and brittleness strongly restrict their processing as solid electrolytes. On the contrary, sulfide-based lithium ion conducting compounds appear to be a much more favorable compound class for solid electrolytes.8 They can be easily processed and densified by cold pressing due to their ductility (low Young’s moduli).9 Furthermore, sulfide-based solid electrolytes show superior conductivities up to 10−2 S cm−1 comparable to, or even beyond those of liquid electrolytes. The most prominent materials are glasses, glass-ceramics and crystalline phases within the quasi-binary Li2S-P2S5 system or within quasi-

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ternary systems with e.g. SiS2, GeS2, Li3PO4, or other additives such as halides.10-13 Outstanding examples are the argyrodite-type phases Li6PS5X (X=Cl, Br, I),14 thio-LISICONtype phases (e.g. Li4-xGe1-xPxS4),15 Li7P3S1116-18 as well as Li10GeP2S1219 and its derivatives20,21 such as 22 Li9.54Si1.74P1.44S11.7Cl0.3, that exhibits the highest Li+ ion conductivity with 2.5  10−2 S cm−1 ever measured in the solid state to date. One fundamental reason often discussed as prerequisite for conductivities that high in sulfide-based systems (besides e.g. Li+ concentration or structural conditions) is the high polarizability of the framework structure.9,23 This concept has already been pursued early by incorporating further soft Lewis-basic anions like iodide ions into thiophosphates. In the early 1980s, first reports indicated quite high conductivities of around 10−3 S cm−1 at room temperature within the system Li2S/P2S5/LiI. While those former investigations focused on compositions with a Li2S : P2S5 ratio of 2:1 including up to 45 mol% of LiI,24,25 more recent studies took systems with Li2S : P2S5 ratios of 7 : 3 (with up to 20 mol% LiI)26 and 3 : 1 with 35 mol%27 and 33 mol% of LiI (= Li7P2S8I),28 respectively, also with oxide-doping29 into account. Rather high conductivities (1.3  10−3 S cm−1), excellent electrochemical stabilities (both against Li metal and up to 10 V) and good battery performances were pointed out,30 making the Li2S/P2S5/LiI system, like related systems,31 highly promising for robust solid electrolyte materials for solid-state secondary batteries. However, as the majority of materials prepared within the Li2S/P2S5/LiI system were vitreous (apart from Li6PS5I), hardly any detailed structural information was gathered in this system, so far impeding proper structure-properties relationships studies.32,33 In this contribution, we report about the discovery of crystalline Li4PS4I, a superionic conductor in the Li/P/S/I system with a novel structure type. On the basis of this crystalline phase, we can now provide an idea of how structural features appear in this system and thus affect the lithium ion conductivity.

RESULTS and DISCUSSION Synthesis In inorganic thiophosphate chemistry, compounds are mostly synthesized by classic solid-state high temperature reactions. Especially for the use as solid electrolytes lithium thiophosphates are often also synthesized through a mechanochemical approach namely high-energy ball milling.8 With regard to industrial production of solid electrolytes, both methods can be unfavorable in terms of costs or scale-up. For the synthesis of Li4PS4I instead, we employed a solvent-based soft chemistry approach that has been utilized similarly for the synthesis of e.g. Li7P2S8I and Li7P3S11.28,34

3 Li2S + P2S5

DME RT

Li3PS4·DME + LiI DME RT

2 Li3PS4·DME

(1)

200°C Li3PS4·DME·LiI −DME Li4PS4I (2)

In a preceding step of this synthesis, Li3PS4·DME was prepared (Eq. (1)) as the essential starting material by stirring the sulfides Li2S and P2S5 in 1,2-dimethoxyethane (DME). The identification of the compound was carried out by crystal structure solution (CCDC 1514533)35 and thermogravimetric measurements (see Supporting information, Tables S1, S2 and

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Figures S1, S2, S3). Li3PS4·DME, in contrast to the Li2S/P2S5 mixture, is well suited to react with LiI according to Equation (2) to form a precursor that is subsequently heated to moderately elevated temperatures finally to yield Li4PS4I. In the initial reaction the yellow solution of LiI in DME is rapidly decolorized when one mole of poorly soluble Li3PS4·DME is dispersed. Consequently, the originally dissolved LiI seems to get absorbed by the solid and incorporated into the structure of Li3PS4·DME to form a respective single-source precursor with a composition of Li3PS4·DME·LiI (cf. reaction scheme in the Supporting Information, Figure S4). After removing excess solvent in vacuum, further crystal DME molecules get expelled when the precursor is annealed at temperatures in the range of 90-250 °C (the ideal temperature range is indicated by TG/DTA data shown in the Supporting information, Figure S5) and Li4PS4I crystallizes in form of a microcrystalline powder (Figure S6). Deviations from the stoichiometric ratio 1:1 either lead to (unreacted) LiI or β-Li3PS4 impurities (originated from unreacted Li3PS4·DME) after the heat treatment. Crystal Structure Analysis The crystal structure of Li4PS4I was solved ab initio from powder diffraction data. All reflections observed in a preliminary Cu-Kα1 X-ray diffraction pattern could be unambiguously indexed based on a primitive tetragonal unit cell with refined lattice parameters a = 8.484(2) Å and c = 5.937(2) Å. Systematic absences referred to the space groups P4/n (no. 85) or P4/nmm (no. 129). Structure solution succeeded in P4/nmm using the charge-flipping algorithm as implemented in TOPAS Academic36 and resulted in a partial structural model consisting of I, S and P atoms only. The Li positions were located by refinement against neutron powder diffraction (NPD) data and subsequent difference Fourier syntheses. The final structure model was refined against both Mo-Kα1 X-ray diffraction and NPD data in a combined Rietveld refinement (Figure 1a, Tables S3, S4 and Figure S7).37 The presence of one phosphorus site found in the asymmetric unit was confirmed by one single resonance in the 31P solid-state MAS NMR spectrum (Figure S8).

Figure 1. a) Rietveld refinement (observed, calculated diffraction patterns as well as difference profiles) against powder X-ray (MoKα1, 0.7093 Å) and neutron (inset, 2θ=79º−106º detector bank) diffraction data; allowed peak positions are marked by vertical lines: Li4PS4I (blue), LiI (green), Li2S (orange); b) Crystal structure of Li4PS4I showing the arrangement of the PS43−-tetrahedra (blue) and the I− (purple) as well as the Li+ ions (green) along [010].

The crystal structure of Li4PS4I consists of Li+ and I− as well as of PS43− ions. As illustrated in Figure 1b the isolated PS43− tetrahedra are formally arranged in layers perpendicular to the c-axis kept apart by the I− ions. Along [101] the arrangement of the tetrahedra and their orientation look similar to that one in α-Li3PS4,38 the high-temperature polymorph of Li3PS4 (Figure 2a). However, in Li4PS4I – as obvious from the formula –

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every second PS43− tetrahedron is replaced by I− (Li3PS4  Li2(PS4)0.5I0.5 ≡ Li4PS4I). A relationship of the Li4PS4I structure with the lithium argyrodites Li6PS5X that also contain halide ions can be derived as shown in Figure 2b. Starting from the Li6PS5X structure (Li+ and tetrahedral S atoms are omitted) the substitution of I to P and S to I as well as addition of I on the vacant Wyckoff position 4d leads to the CsCl structure type with half the lattice parameter a. A subsequent cell transformation with a √2 lattice parameter expansion and a symmetry descent from Pm3¯m to P4/nmm as well as a c-axial displacement of I results in the unit cell of Li4PS4I with the Pand I-positions on Wyckoff positions 2b and 2c, respectively.

Figure 3. Environments of the Li+-sites in Li4PS4I (Li+ green, PS4tetrahedra blue, I− purple).

Figure 2. a) Comparison of the tetrahedra arrangement of αLi3PS4 with Li4PS4I where half of the tetrahedra are replaced by I− (PS4-tetrahedra blue, I− purple, Li+ ions are omitted for clarity reasons); b) Structural derivation of the Li4PS4I structure from the argyrodite structure type (phosphorus blue, I− purple, sulfur yellow).

The Li+ ions were found to be situated within the voids of the PS4I4− substructure distributed over five crystallographic independent sites. The most occupied Li(1) (sof 0.68(3)), Li(2) (sof 0.58(4)), and Li(4) (sof 0.53(3)) as well as the lower occupied positions Li(3) (sof 0.38(2)) and Li(5) (sof 0.08(2)) sum up to an overall occupancy resulting in an electroneutral formula within the standard deviation (Li4.1(1)PS4I). The environments of the Li+ ions are shown in Figure 3. Li(1) and Li(4) are six-fold coordinated by four sulfur atoms in a planar manner and two axial I− ions. The distances Li···S are 267.4(5) and 279.6(1) pm, respectively, and in the typical range found for those arrangements (e.g. in β-Li3PS4: 240.9-260.7 pm). The Li+ ions at Li(2) are tetrahedrally coordinated by two PS43− ions each with typical Li···S contacts of 241.0(1) pm formally forming chains along [001] consisting of alternating and edgesharing PS4- and LiS4-tetrahedra. Li(3) and Li(5) are five-, and fourfold surrounded by 4+1 and 3+1 sulfur atoms and I− ions, respectively. All interatomic distances, summarized in Table S5, are in the range of the sum of the respective ionic radii,39,40 although the Li(1)···I contact with 261.7(21) pm appears to be rather short (distance in LiI: 274.2 pm). However, taking the partial occupancy and a high mobility into account, it is still in a reasonable range.

In order to get an idea of the Li+ migration pathways in Li4PS4I, the crystal structure has been analyzed regarding voids and channels with the Voronoi-Dirichlet polyhedra concept as implemented in the program package ToposPro.41,42 Anurova et al. identified a considerable number of Li+ containing compounds with infinite diffusion pathways as solid electrolytes or as promising candidates employing this analysis method.43 The Li+ migration pathways calculated for Li4PS4I are depicted in Figure 4a. Passing round the PS43− tetrahedra and the I− ions, the Li+ diffusion in Li4PS4I seems to follow a complex, three-dimensional (3D) pathway including all Li sites. Even Li(2), the Li+ position that seems best defined and slightly isolated, is part of the 3D migration pathway. Although this calculation is only a qualitative estimation, isotropic lithium ion conductivity is consequently suggested in Li4PS4I.

Figure 4. a) Possible Li+ migration pathways in Li4PS4I suggested by calculations of the voids with ToposPro (here the structure is shown in origin choice 1 for clarity reasons); b) Arrhenius plot for the total lithium conductivity of Li4PS4I exhibiting an activation barrier of 0.43 eV;44 the conductivities were derived by impedance spectroscopy measurements (inset).

Li+ Mobility Analysis The Li+ ion conductivity of Li4PS4I was studied utilizing impedance spectroscopy. Impedance spectra in Nyquist presentation exhibit a partly resolved semicircle in the high frequency region and a steep tail in the lower frequency region (inset in Figure 4b and Figure S9a). As the first semi-circle could only clearly be resolved up to 2 MHz, bulk and grain boundary transport could not be further distinguished. The observed total lithium ion conductivity shows an Arrhenius behavior (σ·T vs. T−1) for the samples that leads to activation energies of EA between 0.37 eV (35 kJ mol−1) and 0.43 eV

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(41 kJ mol−1). The measured room temperature conductivities range between 1.210−4 and 6.410−5 Scm−1 and are remarkably higher than the ones of the phases Li4PS4I is formally composed of, i.e. Li3PS4 (β: 910−7 Scm−1, γ: 310−7 Scm−1)45,46 and LiI (~10−7),47 but still lower than of β-Li3PS4 that was prepared from the precursor Li3PS4  3THF (210−4 Scm−1)48 or the glassy Li/P/S compositions mentioned above. Although the steep tail in the lower frequency region of the Nyquist plot already suggests predominantly ionic conduction, polarization measurements were performed and the electronic transference number of 4.8(5)10−6 was evaluated, which proves the solid electrolyte character of this lithium thiophosphate (Figure S9b). For a deeper insight into the Li+ ion diffusion process in Li4PS4I temperature-dependent 7Li nuclear magnetic resonance (NMR) spectra were recorded on a static sample at a resonance frequency (ω0/2π) of 77.8 MHz at temperatures between 250 K and 540 K (Figure 5a). At 250 K a broad line is visible at 0 ppm with a width of about 14 ppm (1.1 kHz) representing the so-called central transition (|+1/2〉 → |-1/2〉) of the Li nuclei (nuclear spin I(7Li) = 3/2). Furthermore, a broad contribution is visible in the range from +50 to -50 ppm representing the quadrupolar satellite transitions (|+3/2〉 → |+1/2〉 and |-1/2〉 → |-3/2〉). The fact that this broad contribution does not show any substructure can be explained either by the presence of structural disorder or by a fast motion of the Li+ ions already at these temperatures. When the temperature is increased to 540 K, the quadrupolar contribution disappears and the central line reveals a continuous narrowing with a final linewidth of 4 ppm (300 Hz). This reflects the onset of Li+ diffusion with hopping rates well above the low-temperature linewidth of 1 kHz.49,50 Furthermore, the central line shows a continuous shift from 0 ppm at 250 K to −3 ppm at 540 K. The full width at half maximum (FWHM) of the central transition vs temperature is displayed in Figure 5b. It can be seen that most of the motional narrowing is already completed at 320 K, while continuous narrowing still occurs above this temperature with the absence of a clear plateau. This might be a hint for a complex Li diffusion mechanism involving multiple Li sites that are present in the structure. The spin-lattice relaxation rates T1−1 of the 7Li nuclei were measured in the same temperature range and are plotted vs inverse temperature in Figure 5c. At low temperatures the relaxation rate shows a continuous increase from values below 1 s−1 at temperatures below 320 K to values around 5 s−1 at 500 K. At 500 K the relaxation rate passes a clear maximum characteristic of a diffusion induced spine-lattice relaxation. This occurs exactly when the hopping rate of the Li+ ions reaches the Larmor frequency of their nuclei (τ−1 ≈ ωL). At temperatures above 540 K the relaxation rate then decreases monotonically, while we took care not to go beyond the decomposition temperature of the sample (see TG data in Figure S10). In this temperature range all measurements were highly reproducible during several heating and cooling scans and the magnetization transients could be well described with a mono-exponential function. The behavior of T1−1 vs T−1 above 410 K can be well described with a socalled BPP behavior (dashed line in Figure 5c).51,52 From this behavior we extracted microscopic diffusion parameters such as the hopping rate of the Li+ ions (τ−1 = 5·108 s−1 at 500 K) and the activation barrier for single Li+ jumps (EA = 0.23(1) eV). From these values, using the Einstein-Smoluchowski equation53,54 with an average jump length estimated from the

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shortest Li-Li distance (≈ 1.5 Å), a diffusion coefficient of 2·10−12 m2/s is estimated (at 500 K). This value is converted, using the Nernst-Einstein equation and a Li+ concentration nc of 1.871028 m−3, to a bulk Li+ ion conductivity of 1.3 · 10–3 S·cm–1 at 500 K. It is remarkable that the T1−1 values below 410 K show a distinct shoulder which might hint at a complex diffusion mechanism involving multiple Li sites and even faster hopping on shorter length scales. Comparing the results from the NMR relaxometry experiments with the impedance measurements, the σ-values extracted from the T1 measurements came out to be slightly lower than the extrapolated value from the Arrhenius plot of the impedance data. However, it might be the case that the conductivities measured at low temperatures are more influenced by grain boundaries (high EA) than the values observed at higher temperatures where the inner-gain (bulk) contribution (low EA), which is exclusively detected by NMR, dominates. Such a non-Arrhenius behavior with a flattening curve towards high temperatures was also observed for the argyrodite phase Ag9FeS4.1Te1.9.55

Figure 5. a) The static 7Li NMR spectra of Li4PS4I for temperatures between 250 K and 540 K. b) The corresponding linewidths. − c) The 7Li NMR spin-lattice relaxation rate T1 1 in the same temperature range.

CONCLUSIONS With Li4PS4I, we identified one of the first crystalline thiophosphates in the quasi-ternary system Li2S–P2S5–LiI that is considered as origin for robust solid electrolyte materials in solid-state secondary batteries systems. Although it seems that this crystalline phase was observed already in earlier reports,26,27,28,56 deep structural insight has not been gained until now. Li4PS4I exhibits a new structure type with a layer-like arrangement of isolated PS43− tetrahedra in which Li+ ions diffuse in a complex 3D migration pathway system with a lithium ion conductivity of around 1.210−4 Scm−1 at room temperature. With the structural basis, now, the system can be further optimized in terms of the transport properties (e.g. Li+ ion conductivity) required in all-solid-state batteries. Further benefits of the compound emerge from its synthesis approach. The solvent-based method is not only cost-efficient and easy to upscale, but gives the opportunity to tackle one of the major challenges for all solid state batteries, the formation of proper

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active material/solid electrolyte interfaces.8,31,57 With e.g. the cathode active material present in the reaction mixture during the synthesis process, the creation of an intimate contact to Li4PS4I as the solid electrolyte might be feasible and so a well suitable composite cathode for all solid state battery systems. In addition, it should be possible to prepare thin films of Li4PS4I from precursor solutions at relatively mild thermal conditions, offering new paths for the deposition of ionconducting separator films in solid state batteries.

ASSOCIATED CONTENT Supporting Information Experimental details (Syntheses, X-ray and neutron diffraction, NMR, impedance spectroscopy, TG/DTA, SEM/EDX) TG/DTA data, Crystallographic data of Li3PS4·DME and Li4PS4I, Reaction scheme, SEM images, 31P NMR data, Conductivity and polarization data of Li4PS4I. The Supporting Information is available free of charge on the ACS Publications website. SupportingInformation_Li4PS4I.pdf

AUTHOR INFORMATION Corresponding Authors * Dr. S. J. Sedlmaier, Prof. Dr. J. Janek, BELLA Battery and Electrochemistry Laboratory Institute of Nanotechnology (INT) E-mail: [email protected], [email protected]

Present Addresses Dr. Stefan J. Sedlmaier BMW group, Research Battery Technology, Petuelring 130, 80788 Munich, Germany

Author Contributions All authors have given approval to the final version of the manuscript.

Funding Sources This study is part of the projects being funded within the BASF International Network for Batteries and Electrochemistry.

Notes The authors declare no competing financial interests.

ACKNOWLEDGMENT We are grateful to the ISIS facility including the GEM Xpress service for access to the neutron instrument. This study is part of the projects being funded within the BASF International Network for Batteries and Electrochemistry.

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