Lithium Ion Conductivity in Double Anti-perovskite Li6.5OS1.5I1.5

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Lithium Ion Conductivity in Double Anti-perovskite Li6.5OS1.5I1.5: Alloying and Boundary Effects Hongjie Xu, Minjie Xuan, Weidong Xiao, Yonglong Shen, Zhenzhen Li, Zhuo Wang, Junhua Hu, and Guosheng Shao ACS Appl. Energy Mater., Just Accepted Manuscript • DOI: 10.1021/acsaem.9b00861 • Publication Date (Web): 22 Jul 2019 Downloaded from pubs.acs.org on July 22, 2019

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ACS Applied Energy Materials

Lithium Ion Conductivity in Double Antiperovskite Li6.5OS1.5I1.5: Alloying and Boundary Effects Hongjie Xu1,2, Minjie Xuan1,2, Weidong Xiao1,2, Yonglong Shen1,2, Zhenzhen Li1,2, Zhuo Wang1,2*, Junhua Hu1,2*, Guosheng Shao1,2* 1State

Center for International Cooperation on Designer Low-Carbon & Environmental

Materials (CDLCEM), Zhengzhou University, 100 Kexue Avenue, Zhengzhou 450001, China. 2Zhengzhou

Materials Genome Institute, Building 2, Zhongyuanzhigu, Xingyang

450100, China. E-mail: [email protected]; [email protected]; [email protected] Keywords ：

Solid electrolyte, double anti-perovskite, ionic conductivity, grain

boundary, surficial reconstruction

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ABSTRACT Solid electrolytes based on theoretically identified double anti-perovskite phases Li6OSI2 were successfully synthesized. Experimental characterization supported theoretical prediction that S substitution of O leads to stabilization of the double antiperovskite structure and lattice softening to significantly enhance ionic conductivity, so that the total Li+ conductivity in Li6.5OS1.5I1.5 was two to three orders better than that of the best stoichiometric anti-perovskite phase Li3OCl. However, both anti-perovskite and double anti-perovskite materials are fundamentally susceptible to surface reconstruction, which is behind significant boundary resistances typically known for materials based on anti-perovskite hali-chalcogenides. Such a surface related problem was then effectively reduced through amorphous phase formation, thus offering a feasible route to exploit the full potential of this class of new materials as competitive candidates for solid Li-ion batteries.


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Introduction Lithium ion batteries have been extensively exploited in recent decades, owing to ever increasing demands in energy storage particularly for powering portable electronic devices and electric vehicles. However, safety is a major concern for current commercial technologies, due to high flammability of liquid organic electrolytes.1-4 Significant efforts have therefore been directed towards developing solid-state electrolytes (SSE), in order to deliver all-solid-state batteries (ASSB) that are fundamentally safer.5-12 The first major landmark breakthrough in developing SSE was realized through achieving superb ionic conductivity in the LGPS family of materials, typically Li10GeP2S1213-14 and Li9.54Si1.74P1.44S11.7Cl0.3.15-16 However, the LGPS materials are not electrochemically compatible with Li anode,17 thus hindering its use with Li anode for the highest capacity of 3860 mAh g-1 and lowest chemical potential of -3.04 V vs. the standard hydrogen electrode.18 Recently, a new class of SSE materials has been developed under the guidance of systematic modelling in tuning materials chemistry in argyrodite based on cubic hali-chalcogenides, resulting in significantly lowered activation energy for Li+ transportation with respect to the LGPS family (0.17 eV vs. 0.22 – 0.25 eV in LGPS), together with highly improved electrochemical stability with Li anode.8, 10-12 Great attention have also been paid to Li-rich anti-perovskite hali-oxides Li3OX (X = halogen anions Cl, Br, etc.), owing to its being more electrochemically compatible with the lithium anode.19-20 However, it has been found that the ionic conductivity in Li3OX is insufficient. Li+ conductivity realized in Li3OCl was only about of 0.85 mScm1

at room-temperature, with an activation energy Ea=0.26 eV. While it was reported that

mixing in the halogen site such as Li3OCl0.5Br0.5 led to slightly higher ionic conductivity of 1.94 mScm-1,21 one notes that huge additional resistance due to double layer of charge distribution was attributed to “grain boundary” resistance, being over 200 Ω at 3 ACS Paragon Plus Environment


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a rather high temperature of 190 ℃.22 Indeed, such reported “grain boundary” resistance was orders higher than the “bulk” ion conductivity, such that the effective total ionic conductivity was rather poor, thus preventing their application in high-performance solid batteries.23 Recently, through systematic first-principles simulation, we have formulated a new class of double anti-perovskite electrolytes based on Li6OSI2, which is expected to deliver superb ionic conductivity together with very small activation energy for Li+ diffusion.10 Here in this work, we report experimental efforts in synthesis and characterization of this new class of materials, with particular efforts in verifying the theoretically predicted remarkable S-effect on chemical ordering and hence highly improved Li+ conductivity, with respect to that of the best anti-perovskite hali-oxide Li3OCl. We confirm that the ionic conductivity can be significantly improved via the formation of the double anti-perovskite structure, while there is still usual concern of significant so-called “grain boundary resistance”, as has been well recognized in Li3OCl. First-principle calculations show that such a boundary problem is attributable to easy surficial reconstruction in this class of materials, being qualitatively in accord with XPS evidence. Efforts in reducing such surficial segregation through amorphous phase formation leads to further enhancement of ionic transportation, even though there still exists boundary resistance due to inadequate bonding between amorphous phase particles. The insightful outcome of the current work thus provides useful guidance in realization of the full potential of double anti-perovskite as an attractive solid-state electrolyte for safe Li-ion batteries. Experimental Methods Materials synthesis Crystalline materials for Li6OSI2 and Li6.5OS1.5I1.5 were synthesized by ball milling. The whole processes were carried out in a glovebox under Ar atmosphere to isolate the materials from humidity and air. The multi-constituent electrolytes were 4 ACS Paragon Plus Environment

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produced through the following steps: (a) ball milling of mixed powders of Li2S, Li2O and LiI (all from Aladdin, 99.99%) with designated molar ratios at 550 rpm for 8 h to obtain precursor mixtures of powders. LiI peaks were present in the XRD pattern shown in Figure S1(b); (b) thermal annealing of the precursor powder at 190 ℃ for 10 hours, and then cooling down to room temperature at 5K per minute. As shown in Figure S1(b), some weak peaks from double anti-perovskite (indicated by arrows) emerged; (c) ball milling at 350 rpm for 4 h; (d) pressing milled powder into pellets under 500 Mpa, followed by annealing at 220 ℃ for 15 h, and then cooled down to room temperature by 5K per minute. Li3OCl ceramic materials was made in a similar way, for which Li2O and LiCl were milled by the molar ratio of 1:1, followed by annealing the pressed pellets at 250 ℃ for 10 h, and then cooling quickly down to room temperature. Amorphous samples of Li6OSI2 was synthesized by mixing LiI, Li2O and Li2S in a molar ratio of 2:1:1. The mixed precursor with the weight of 9g were introduced into 40 ml solution of anhydrous 1,2 dimethoxyethane (DME Aladdin 99.5%). The mixture was stirred at room temperature in a bull bottle for a week. Then, the colorless precursor was filtered out and dried in vacuum at 60℃ for 24h to remove residual DME. The solid precursor was then annealed 230°C in a tungsten crucible for 10 h to obtain a noncrystalline powder product. The powder was finally pressed into 0.25 g pellets under 500 Mpa. Materials characterization X-ray diffraction (XRD, Rigaku Ultima IV) with Cu-Ka radiation was employed for identification of phases, with samples being sealed within polyimide film in the glovebox for protection from ambient atmosphere. Field emission scanning electron microscopy (SEM, ZEISS, SIGMA 500/VP) was used for morphological and microstructural characterization.



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X-ray photo-electron spectroscopy (XPS, VG Mulitilab 2000) with the Al Kα xray source (1486.6 eV) was adopted to characterize surficial alloy chemistry. Ar ion etching of the surfaces was followed to characterize the materials below the exposed surfaces (bulk). The melting and crystallization temperatures of the composites were determined via differential scanning calorimetry analysis (DSC, DSC214, Netzsch), within a closed alumina crucible at a heating rate of 10 K min-1 from room temperature to 500 ℃ under an Ar flow for protection from the ambient atmosphere. Ionic conductivity measurement Ionic conductivity was measured by electrochemical impedance spectroscopy (EIS) with an applied frequency from 100 Hz to 7 MHz, using a Schlumberger Solartron 1260 frequency response analyzer at a sinusoidal amplitude of 10 mV. The samples were pressed into pellets and heated in vacuum at the required temperature for each composition. Both sides of the pellet (diameter 15 mm; thickness 0.48mm) were coated with Au to act as current collectors. The pellet was then put into a copper clapping assembly within Ar filled glove box, which helped secure good contact to electrodes and isolate the sample from air. The clapped sample was then subject to any specified temperature in a pre-dried testing chamber, keeping for 0.5 h to settle the sample temperature before impendence measurement. Calculation method First principles calculations are carried out in the framework of the density functional theory (DFT), on the basis of extensively tested approaches.8-12, 24-26 Details for the calculations are presented in the Supplementary Information. Results and Discussion Materials characterization 6 ACS Paragon Plus Environment

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The general approach for materials synthesis involves ball milling of mixtures of constituent compounds followed by subsequent annealing. The overall reaction for the formation of the stoichiometric compound of the double anti-perovskite, for example, is

𝑳𝒊𝟐𝑶 + 𝑳𝒊𝟐𝑺 +𝟐𝑳𝒊𝑰→𝑳𝒊𝟔𝑶𝑺𝑰𝟐. The heating and cooling DSC traces from ball-milled Li6OSI2 powders are shown

in Figure 1. The presence of a single-endothermic peak in the heating trace is indicative of phase transition at 193.4 ℃. The peak at 381.6 ℃ corresponds to the liquidus of the alloy. In the cooling process, the first exothermic valley is the solidus of the double anti-perovskite phase, which is consistent with the liquidus but at slightly lower temperature due to thermal shooting during both heating and cooling. The crystallization temperature in the cooling curve for the material is shown to be 176.1 ⁰C. It is about 20% higher than the theoretically predicted temperature of 145 ⁰C for the decomposition of the double anti-perovskite structure into binary species at lowered temperature, which is considered reasonable given that only the phonon entropy was included in the DFT assessment of free energies.10 The endothermal peak in the heating curve at 193.4 oC is attributable to disappearance of binary phases in the as-milled powder. On the basis of thermodynamic modelling,10 the double anti-perovskite phase is stable between the melting and the highest crystallization temperature in the DSC traces. Such DSC feedback was thus combined with the theoretical outcome to optimize the processes for material synthesis, as is summarized in the experimental session. Besides, the DSC traces for slightly off-stoichiometric compounds are consistent with that of Li6OSI2. Therefore, similar processes for material synthesis of off-stoichiometric compounds were adopted. XRD patterns from three alloys, Li6.5OS1.5I1.5, Li6.25OS1.25I1.75 and Li6OSI2 are presented in Figure 2. The pattern of Li6OSI2 agrees with the calculated pattern using lattice parameters from the theoretical work,10 and off-stoichiometric compounds maintained the main characteristics of Li6OSI2. The XRD pattern of the anti-perovskite Li3OCl, on the other hand, is rather different. The symmetry of Li6OSI2 is 𝐹𝑀3 7 ACS Paragon Plus Environment


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𝑀(225) with a lattice parameter of 8.578 Å, while that of Li3OCl is 𝑃𝑀3𝑀(221) with a lattice parameter of 3.910 Å.10 The structural characteristics for the double antiperovskite and anti-perovskite phases are shown in Figure 2(c, d). A closer look at the patterns, as shown in Figure 2(b), reveals some peak shift towards

lower diffraction

angles due to off stoichiometry from Li6OSI2, indicating lattice expansion in the sequence of Li6OSI2 < Li6.25OS1.25I1.75 < Li6.5OS1.5I1.5. The lattice expansion out of offstoichiometric compositional deviation is another factor for lattice softening, in addition to the S effect. This offers further evidence why the combined offstoichiometric and S effect leads to radical improvement for Li+ conductivity.8, 11 The relative densities for different pellets are similar, being 87.4% for Li3OCl, 84.65% for Li6OSI2 and 88.8% for Li6.5OS1.5I1.5. Figure 3 shows typical SEM morphologies after sintering, taking the Li6OSI2 alloy as an example. Traces of particulate powder features can be observed on the surface of the annealed pellet. The fractogram of the pellet demonstrates dendritic features from microcrystals (arrow indicated), suggesting intergranular and inter-particular fractures due to weaker bonding between grains and particles alike. Overall both particulate and grain boundaries exist, together with some associated porosities in the pellets. Impedance spectroscopy Electrochemical impedance spectroscopy was carried out for samples of Li3OCl and Li6.5OS1.5I1.5. Figure 4(d) and Figure 5 show EIS spectra at different temperatures from 75 to 115 ℃, with insets showing equivalent circuits from numerical fitting using the Z-View code for impedance analysis. Further details are shown in enlarged spectra in Figure S2. Each Nyquist plot of EIS spectrum typically consists of a semicircle and an upward linear tail. The diameter of a semicircle is an electrochemical measure of space charge region due to a double layer of polarized charges, leading to a constant phase element (CPE1) that is parallelly connected with a resistance element (R1) due to difficulty in charge transfer. The charge transfer resistance R1 term was widely termed as “grain boundary” resistance in solid electrolytes,14, 18 especially those based on the 8 ACS Paragon Plus Environment

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anti-perovskite of Li3OCl,11, 27-28 while an ohmic resistor in series (R0) was generally attributed to transport limited resistance owing to the “bulk” materials.28-31 It should be pointed out, however, that the presence of a CPE is not necessarily associated with grain boundaries, since polycrystalline materials with excellent ionic conductivity are free of such a CPE element in the equivalent circuits11 with grain boundaries being widely recognized as fast diffusion channels in condensed matter.32 Therefore, there is little physical basis for the existence of significant “grain boundary” resistance in solid electrolytes, unless passivated boundary layers are present. This could arise either from detrimental chemical segregation at grain boundaries or from poorly bonded particles. Indeed, as was demonstrated recently by Sun et al., the CPE disappeared in EIS and thus leading to significantly improved ionic conductivity when inter-particular/intergranular bonding was improved in an anti-perovskite compound of Na3OBF4 with hot pressing.33 The total effective resistance for a solid electrolyte should be the sum of R0 and R1. As presented in Figure 4d, the R1 term of Li3OCl is extremely large from 74160 to 4597 Ω at different temperatures ranging from 75 to115 ℃ (Detail data for Li3OCl was presented in Table S1), it has been considered as the main concern for this class of solid electrolytes. The fitted values for each component in the equivalent circuits of EIS plots from the double anti-perovskite Li6.5OS1.5I1.5 and anti-perovskite Li6OSI2 alloys are listed in Tables S2 and S3, respectively. In detail, the boundary resistance of Li6.5OS1.5I1.5 decreased from 1180 to 232.8 Ω with the temperature being increased from 75 to 115 oC,

while corresponding R1 for the anti-perovskite Li3OCl were 74160 and 4597 Ω.

Comparing the like with like, the S effect in the double anti-perovskite phase led to dramatically reduced R1 for about 20 times, from 4597 down to 232.8 Ω at 75 oC. On the other hand, R0 of Li6.5OS1.5I1.5 experienced little change with increasing temperature (Table S2), while R0 of Li6OSI2 (Table S3) even increased with increasing temperature. This indicates that R0 could not be attributed to “bulk” resistance alone for an electrically insulating materials such as solid electrolytes, or it would be expected to 9 ACS Paragon Plus Environment


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decrease with increasing temperature instead. The R0 term is therefore more likely to be owing to imperfect ohmic contact between the electrolyte and the metal contact, for example, Schottky junction due to mismatching work function of the metallic electrode and the electrolytes of interest. Similar phenomena for R0 vs. T were also observed in other solid electrolytes such as LGPS 14 and garnet.18 The temperature dependent Li+ conductivity from these two alloys are plotted in Figure 6 (a), using the total resistance data. One can see from the mean square displacements shown in Figure 6(b, c) that only Li+ ions experienced long-range transportation. The two experimental lines with log (𝜎𝑇) vs. 1000/𝑇 from Li6.5OS1.5I1.5 and Li3OCl compounds are below AIMD predicted lines for bulk singlecrystals. In addition to poorer total conductivities, the experimentally derived activation energies are also considerably larger than the theoretically predicted values. The total Li+ conductivity ( ) for Li3OCl measured from experiment was only 3.66×10-4 mS cm-1 even at 75 ℃, with a huge activation barrier of Ea =0.874 eV as indicated in Figure 6(a), which is much worse than the theoretically predicted value (0.27 eV for interstitial Li+ and 0.31 eV for vacancy diffusion mechanism).10, 34,27, 35 This suggests that particle boundaries do seem to be a main issue for the anti-perovskite materials. Such a particle boundary problem also existed in the double anti-perovskite compounds, so that the experimentally derived activation energy is also significantly bigger than the theoretically predicted data (0.48 eV vs. 0.18 eV). Comparing with the anti-perovskite phase, there is around two orders of enhancement of the overall conductivity owing to S effect in the double anti-perovskite phase (2.28×10-2 mS cm-1 at 75 ℃ in Li6.5OS1.5I1.5). Surface chemistry Boundary resistance from ionic transportation has been widely accepted to be a major issue in solid electrolytes based on anti-perovskite hali-chalcogenides, it is necessary to probe for the fundamental reasons. Dawson et al. studied the grain 10 ACS Paragon Plus Environment

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boundary effect using classic molecular dynamics simulation over some artificially constructed grain boundaries. 23, 36-37Here we make further effort using first-principles modelling on low energy surfaces which dictate particle boundary formation, in order to elaborate the role of surface chemistry behind such boundary problems. We start with analyzing the structural reconstruction over the closely pact (001) surface of Li3OCl and Li6OSI2 respectively, since they tend to be the exposed surfaces or at large angle grain boundaries. Convergence tests of surface energies are carried out over 2 to 8 layers of octahedrons in the [001] direction, Figure S3 (a, c). Based on the surface energy tests with the increase of layers as displayed in Figure S3 (b, d), the surface energy for Li6OSI2 becomes convergent at six layers of octahedrons, while the surface energy for Li3OCl converges at seven layers. The slab models with adequate thicknesses are therefore decided, as shown in Figure 7 (a, c) for the two phases. We then use USPEX as a powerful material genome tool to search for globally stable configuration as reconstructed surficial structures, with two bottom layers of octahedrons fixed, Figure 7(a, c). Chemical species in other layers are set to be completely free from forming buffer or surficial regions.38 The identified surficial structures were then subjected to further geometric relaxation, with the energy of the fully-relaxed surficial structure for Li3OCl, Figure 7b, being 0.00446 eV/atom lower than that of the initial structure, Figure 7(a). For the double anti-perovskite Li6OSI2, surficial reconstruction shown in Figure 7(d) is driven by 0.0187 eV/atom, which is more than 4 times of that for the anti-perovskite phase. This suggests that the surficial and associated boundary segregation would be more likely in the latter, which could lead to difficulty in transportation of Li+ across particle/grain boundaries. The corresponding changes in partial density of states (PDOS) are shown in Figure 8. In the case of Li3OCl, surficial reconstruction leads to broadening of PDOS for each element, largely towards lowered energies, Figure 8(a, c). The most pronounced feature lies in the weight center for the PDOS of O being shifted about 0.3 eV downwards, thus enhancing overall binding to Li+ whose peak coincides with that 11 ACS Paragon Plus Environment


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of O-2. For the double anti-perovskite Li6OSI2 phase, the change in PDOS due to surficial restructuring is more dramatic, with peaks for Li+ cations and all anions being shifted downwards. This is suggestive of more significant enhancement of binding of Li+ to anions, which in turn induces electrostatic hindrance to the migration of Li+ ions.8-12, 22, 39 Pair correlation functions (PCF) are calculated to assess the changes of chemical bond lengths after surficial restructuring. For the anti-perovskite Li3OCl phase, the width of the well-defined Li-Cl peak is broadened considerably as displayed in Figure S4(a), though insignificant changes in weight centers of the PCFs are indicative of limited surficial restructuring. Meanwhile, the lengths of Li-O bonds experience little change as shown in Figure S4b. The peak positions do not change for either Li-O or LiCl. On the other hand, significant changes in PCFs are induced by surficial restructuring in the double anti-perovskite phase Li6OSI2 as exhibited in Figure S4(c-e). Apparently, the weight of centers for the PCFs of Li-I, Li-O and Li-S and as well as the peak positions for the latter two are shifted to lower values. Bigger effect in lowered Li+ and anion PCFs in the double anti-perovskite phase is consistent with much larger energetic driving force, as is also demonstrated in more significant changes in PDOS. This suggests that surficial/interfacial Li+ ions are even less mobile with respect to these within the bulk crystals in the double anti-perovskite phase. XPS analysis was carried out to collect spectra from free surfaces and then after surfaces were sputtered in situ with Ar ions for 300 seconds. This provided a qualitative assessment of change in surficial chemistry with respect to interior structure. The guide lines in Figure S5(a) show that the surface of Li3OCl contains more O-2 anions than the inside of the crystals. This agrees with the PDOSs that surficial restructuring leads to enhanced Li-O binding. In the case of the Li6OSI2 sample, the surface enrichment of O-2 anions was even more evident, suggesting even stronger binding to Li+ ions. Overall, the surface analysis results are consistent with modelling results. The surficial


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restructuring leads to rather significant enhancement of binding of Li+ ions to anions, thus leading to greater difficulty in Li+ transportation across particle boundaries. It is worth noting that even though the surface problem for the double antiperovskite phase is more pronounced than that for the anti-perovskite, the overall double layer effect in the former is found to be less serious. This is because the innate great improvement in ionic conductivity in the double anti-perovskite makes it easier for charge transfer, as was shown in recent work on the argyrodite based system.11 This is demonstrated in Figure 9(a), that the total ionic conductivity of the current double anti-perovskite material at 85 oC is similar to the best reported data on anti-perovskite Li3OCl0.5Br0.5 at a temperature more than 100 degrees higher.22 Since the fundamental origin for surficial restructuring causes detrimental boundary segregation, one can envisage that the full potential of the bulk solid electrolyte can be better realized by avoiding particle/grain boundaries. One option is making use of amorphous (am) phase, so that the numerous vacancies and vacancy clusters within the materials can help reduce energetic driving force for surface reconfiguration. The amorphous route was shown to be helpful for effective reduction of the grain boundary problems in anti-perovskites based on Li3OCl.40-44 As is shown in Fig. 9 (b), no grain-boundary resistance was detected in amorphous Li6OSI2 at 85 ℃, with an impressive ionic conductivity being 6.15 mS/cm, which is higher than that of Li3-0.01Ba0.005OCl (4.91 mS/cm).41 It is shown in the inset SEM image that particle boundaries still existed in the amorphous material, Figure 9(b). This suggests further scope for improvement in ionic conductivities, once particle boundaries are removed in the amorphous material. This can be realized by thin-film coating methods such as pulsed laser deposition (PLD), physical vapor deposition and sputtering. In order to evaluate the electrochemical stability of the current electrolytes with metallic anodes, we carried out voltammetry and cycling stability tests using Li/ Li6OSI2(am)/stainless-steel half cells and Li/Li6OSI2(am)/Li symmetric cells respectively. As displayed in Figure S6(a), the absence of any other peaks in the entire 13 ACS Paragon Plus Environment


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voltage range up to 5 V indicates that the amorphous solid-state electrolyte from this work was reasonably stable over a wide voltage range versus metallic Li.

Also, the

cycle stability test using symmetric Li/Li6OSI2(am)/Li cell was carried out by applying a constant current density of 20 µA cm-2 under alternating polarity, as shown in Figure S6(b). There was little change of voltage over a prolonged period of 400 h. Polarization became more serious at a higher current of 40 µA cm-2, Fig. S7, indicating formation of interfacial solid inter-phase. Conclusion Double anti-perovskite Li6OSI2 compound were successfully synthesized under the guidance of recent theoretical work, which predicted superb ionic transportation in this new material system, together with significant thermodynamic and electrochemical stability. Total Li+ conductivity in the alloy with slight Li excess, Li6.5OS1.5I1.5, was found to deliver about two to three orders of improvement of ionic conductivity over the bestperforming stoichiometric anti-perovskite phase Li3OCl. Both anti-perovskite and double anti-perovskite materials are fundamentally susceptible to surface reconstruction, leading to significant boundary resistances. Utilization of amorphous structures are shown to be effective in realizing the full potential of the new compound as a competitive electrolyte system for solid Li-ion batteries.

Associated Content The manuscript is accompanied by Supporting Information containing: -

Computational methods


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Determination of ionic conductivity from EIS data

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Numerically fitted data for elements in the equivalent circuit of the EIS from Li3OCl , Li6.5OS1.5I1.5 and Li6OSI2 at various temperatures

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Phase diagram for Li6OSI2

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Impedance spectra over a broad frequency range from 100 Hz to 7 MHz for samples of Li6OSI2 (75 to 115 ⁰C), with enlarged impedance spectra for Li3OCl, Li6.5OS1.5I1.5, and amorphous Li6OSI2

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Convergence tests on surface energy from 2 to 8 layers of octahedrons in the slabs of Li3OCl and Li6OSI2

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Pair correlation functions (PCF) in Li3OCl and Li6OSI2

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XPS spectra from Li3OCl and Li6OSI2

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The Cyclic voltammograms curve of crystal Li6.5OS1.5I1.5 and the stable of the electrolytes up to 5V

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Galvanostatic cycling of Li/ Li6OSI2(am)/Li with a change current density of 10 μA/ cm2, and 20 μA/ cm2 and 40 μA/ cm2, respectively

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POSCAR file

Author Information Corresponding Author



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E-mail: [email protected]; [email protected]; [email protected]. Author Contributions The manuscript was written through contributions of all authors. All authors have given approval to the final version of the manuscript.

Acknowledgments This work is supported in part by the 1000 Talents Program of China, the Zhengzhou Materials Genome Institute, and the National Natural Science Foundation of China (No. 51001091, 51571182, 111174256, 91233101, 51602094, 11274100). References (1) Armand, M.; Tarascon, J.-M. Building better Batteries. Nature. 2008, 451, 652657. (2) Dunn, B.; Haresh Kamath; Tarascon, J.-M. Electrical Energy Storage for the Grid: A Battery of Choices. Science. 2011, 334, 928-935. (3) Scrosati, B.; Hassoun, J.; Sun, Y.-K. Lithium-ion Batteries. A Look into the Future. Energy Environ. Sci. 2011, 4, 3287-3295. (4) Wang, Q.; Ping, P.; Zhao, X.; Chu, G.; Sun, J.; Chen, C. Thermal Runaway Caused Fire and Explosion of Lithium ion Battery. J.Power Sources. 2012, 208, 210224. (5) Song, A.-Y.; Xiao, Y.; Turcheniuk, K.; Upadhya, P.; Ramanujapuram, A.; Jim Benson, A. M.; Olguin, M.; Meda, L.; Borodin, O.; Yushin, G. Protons Enhance Conductivities in Lithium Halide Hydroxide/Lithium Oxyhalide Solid Electrolytes by Forming Rotating Hydroxy Groups. Adv. Energy Mater. 2017, 1700971. (6) Asano, T.; Sakai, A.; Ouchi, S.; Sakaida, M.; Miyazaki, A.; Hasegawa, S. Inorganic Solid-State Electrolytes for Lithium Batteries: Mechanisms and Properties Governing Ion Conduction. Adv Mater. 2018, 30 (44), 140-162. (7) Bachman, J. C.; Muy, S.; Grimaud, A.; Chang, H. H.; Pour, N.; Lux, S. F.; Paschos, O.; Maglia, F.; Lupart, S.; Lamp, P.; Giordano, L.; Shao-Horn, Y. Inorganic Solid-State Electrolytes for Lithium Batteries: Mechanisms and Properties Governing Ion Conduction. Chem Rev. 2016, 116 (1), 140-62. (8) Wang, Z.; Shao, G. Theoretical Design of Solid Electrolytes with Superb ionic Conductivity: Alloying Effect on Li+ Transportation in Cubic Li6PA5X Chalcogenides. J. Mater. Chem. A. 2017, 5, 21846-21857. (9) Wang, Z.; Shao, G. High-Capacity Cathodes for Magnesium Lithium Chlorine Tri-ion Batteries Through Chloride Intercalation in Layered MoS2: a Computational Study. J. Mater. Chem. A. 2018, 6, 6830-6839.


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Figures

Figure 1 Differential scanning calorimetry (DSC) traces during heating and cooling of Li6OSI2, at heating/cooling rate of 10 ℃ min-1.


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Figure 2 (a) Theoretically predicted XRD pattern for Li6OSI2 (Cal), compared with experimental patterns from synthesized samples Li3OCl, Li6OSI2, Li6.25OS1.25I1.75, and Li6.5OS1.5I1.5. (b) Enlarged partial patterns from 25 to 35°. (c) Li6OSI2 with double antiperovskite type structure. (d) Li3OCl with anti-perovskite type structure.



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Figure 3 SEM images from pressed pellets of synthesized Li6OSI2 powder samples: (a) surface morphology, and (b) cross-section fractogram of pressed pellets after annealing.


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Figure 4 (a) The diameter and (b) thickness of pellet before coating Au. (c) The sealing clamp assembly used for impedance measurement with Au coated pellet to enhance contact with electrodes. (d) The impedance spectra of Li3OCl over a broad frequency range from 100 Hz to 7 MHz, taken at different temperatures from 75 to 115 ⁰C. The typical equivalent circuit is shown as Au coated electrolyte pellet inset.



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Figure 5 The impedance spectra over broad frequency from 100 Hz to 7 MHz for the synthesized samples of Li6.5OS1.5I1.5, taken at temperatures between 75 and 115 ⁰C. Inset shows typical equivalent circuit.

Figure 6 (a) Arrhenius plots presented as log(σT) vs. 1000/T to compare theoretical and experimental data for Li6.5OS1.5I1.5 and Li3OCl. (b) MSD of Li+, O2-, S2-, and I- at 2000 K. (c) MSD of Li+ for double anti-perovskite from 750 K to 2000 K. 24 ACS Paragon Plus Environment

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Figure 7 (a) Initial and relaxed (b) slab models for the (001) surface of Li3OCl; and (c) initial and relaxed (d) slab models for the (001) surface of Li6OSI2.



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Figure 8 (a, c) Calculated partial densities of states in initial and relaxed slabs of Li3OCl. (b, d) Calculated partial densities of states in initial and relaxed slabs of Li6OSI2.


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Figure 9 EIS from (a) crystalline anti-perovskite Li3OCl0.5Br0.5 and double antiperovskite Li6.5OS1.5I1.5; and from (b) amorphous Li6OSI2 and Li3-2*0.005Ba0.005OCl, with insets of cross-sectional SEM image and XRD patterns for amorphous Li6OSI2.



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For Table of Contents Only


Lithium Ion Conductivity in Double Anti-perovskite Li6.5OS1.5I1.5

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