Synthesis and Electrochemical Properties of I4̅ ... - ACS Publications

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Synthesis and Electrochemical Properties of I4̅-Type Li1+2xZn1−xPS4 Solid Electrolyte Naoki Suzuki,*,† William D. Richards,‡ Yan Wang,‡,§ Lincoln J. Miara,§ Jae Chul Kim,∥ In-Sun Jung,⊥ Tomoyuki Tsujimura,† and Gerbrand Ceder‡,∥,# †

Samsung R&D Institute Japan, Minoh Semba Center Building, Semba Nishi 2-1-11, Minoh, Osaka 562-0036, Japan Department of Materials Science and Engineering, Massachusetts Institute of Technology, 77 Massachusetts Avenue, Cambridge, Massachusetts 02139, United States § Advanced Materials Lab, Samsung Research America, 3 Van de Graaff Drive, Burlington, Massachusetts 01803, United States ∥ Materials Sciences Division, Lawrence Berkeley National Laboratory, Berkeley, California 94720, United States ⊥ Samsung Advanced Institute of Technology, Samsung Electronics Co., Ltd, 130 Samsung-ro, Yeongtong-gu, Suwon-si, Gyeonggi-do 16678, Republic of Korea # Department of Materials Science and Engineering, University of California, Berkeley, California 94720, United States ‡

S Supporting Information *

ABSTRACT: I4̅-type LiZnPS4 solid electrolyte was synthesized starting from Li2S, P2S5, and ZnS, and its room temperature ionic conductivity, σ25°C, was ∼10−8 S/cm. To improve σ25°C, compositions of Li1+2xZn1−xPS4 (x = 0.125, 0.25, 0.375, 0.5, 0.625, 0.75, 0.8, 0.9) were synthesized from mixtures of LiZnPS4 and amorphous Li3PS4 (a-LPS). The samples with x ≤ 0.8 were confirmed by means of XRD and Raman spectroscopy to have the I4̅ structure. The maximum σ25°C was 5.7 × 10−4 S/cm at x = 0.625. NMR measurement revealed that Li1+2xZn1−xPS4 (x ≤ 0.625) is almost a single phase material and not a simple mixture of LiZnPS4 and a-LPS, confirming that its crystal structure is responsible for the improved ionic conductivity. Li2.25Zn0.375PS4 (x = 0.625) was reactive to graphite and indium anodes but stable against a Li4Ti5O12 anode and a 4 V class cathode. We fabricated an all-solid-state battery in the form of graphite|a-LPS|Li2.25Zn0.375PS4|Li2O-ZrO2 coated Li(Ni, Mn, Co)O2 and examined its performance. It showed a first-cycle discharge capacity of ∼144 mAh/g−NMC with a capacity retention of 83% after 100 cycles.

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INTRODUCTION Lithium-ion batteries (LIBs) are currently used to power a wide range of portable devices and, in particular, a number of pure electric and plug-in hybrid vehicles. As the energy density of LIBs increases, concerns about their safety becomes more serious. Conventional LIBs use a liquid electrolyte, which is usually flammable and can ignite as a result of battery failures. To reduce this risk, there have been numerous efforts in recent years to develop all-solid-state Li batteries (ASSBs) that use incombustible solid electrolyte (SE). SEs are mainly classified into two types: oxide-type and sulfide-type. Several oxide SEs have been reported such as perovskite-type solid solutions,1 NASICON-type structures,2−5 and garnet-like structures.6−8 All of them are relatively stable under ambient © 2018 American Chemical Society

conditions, and their ionic conductivities have reached practical levels (10−4−10−3 S/cm). However, these oxide electrolytes generally show a large grain boundary resistance induced by point contact among hard oxide particles, and a sintering process at several hundred to a thousand degrees is usually necessary to improve the ionic conductivity at the grain boundary. This makes oxide SEs difficult to use in practical cells. On the other hand, sulfide SEs have many advantages for use in ASSBs. For example, sulfide SEs can easily be pelletized by Received: September 11, 2017 Revised: March 8, 2018 Published: March 9, 2018 2236

DOI: 10.1021/acs.chemmater.7b03833 Chem. Mater. 2018, 30, 2236−2244

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Chemistry of Materials

structural descriptor, the study also suggested a few lithiumcontaining sulfide compounds whose anion structures are close to bcc, as potential candidates for new solid-state lithium ionic conductors. In this study, we examine LiZnPS4a bcc-type material as suggested by Wang et al.36 The synthesis of stoichiometric LiZnPS4 was previously reported by S. Jörgens et al.,37 who reported its crystal structure as I4̅-type, the same as Zn3(PS4)2. This structure contains two alternating layers of corner sharing sulfur tetrahedra, one layer being half occupied by phosphorus and the other containing an ordered arrangement of Li and Zn atoms. The sulfur sublattice of LiZnPS4 is found to be very close to bcc with an a/c ratio of 0.9 and the sulfur atoms displaced on average only 0.29 Å from their ideal positions.36 The ionic conduction properties of the LiZnPS4 have been studied recently by ab initio molecular dynamics simulation (AIMD) using density functional theory. Richards et al.38 predicted that the stoichiometric LiZnPS4 is almost insulative, but high mobility of Li can be achieved in the nonstoichiometric versions of the compound Li1+2xZn1−xPS4. They illustrated Li1+2xZn1−xPS4 as a solid solution of LiZnPS4 and the Pmn21 phase of γ-Li3PS4 (Li1+2xZn1−xPS4 = (1 − x)LiZnPS4 + xLi3PS4), and produced the binary phase diagram.38 It was shown in the diagram that LiZnPS4 and Li3PS4 are soluble in a wide composition range. In this case, Li1+2xZn1−xPS4 forms the I4̅ structure, in which Zn is partially substituted to Li and excess Li stays in vacant tetrahedral sites in the P layer, which drastically improves the ionic conductivity of stoichiometric LiZnPS4. The compositions Li1+2xZn1−xPS4 are predicted to show room temperature ionic conductivity as large as 50 mS/cm at x = 0.5. In this work, we report the experimental synthesis and electrochemical properties of Li1+2xZn1−xPS4 with various x. We confirm that stoichiometric LiZnPS4 has a low room-temperature conductivity and a high activation energy, and that the conductivity of Li1+2xZn1−xPS4 greatly increases as the lithium concentration increases, in qualitative agreement with firstprinciples calculations.38 We also discuss possible reasons why the very high predicted conductivity could not be achieved in these experiments.

cold-pressing at room temperature. These pellets usually do not show a large grain boundary resistance. Conductivity can be further improved by hot-pressing or spark-plasma sintering.9,10 Easy production is a strong point in the manufacturing process. There are some particular difficulties in the development of sulfide based ASSBs. One is the large interfacial resistivity between the SE and cathode, which is believed to be induced through the mutual diffusion of component atoms of SE and cathode, such as sulfur, phosphorus, oxygen, and transition metals leading to decomposition products forming at the interface.11,12 N. Ohta et al.,13 however, circumvented this problem by coating the surface of the LiCoO2 cathode particle with Li4Ti5O12, and found that the power density was dramatically increased, comparable to commercial lithium-ion batteries with organic-solvent liquid electrolytes. This has spurred a vast amount of studies about the improvement of the barrier layer on the cathode materials.11−20 Recently, Tsujimura et al.21 investigated the effect of the coating layer in more detail using Time-Of-Flight Secondary Ion Mass Spectroscopy (TOFSIMS) and X-ray Absorption Near Edge Structure (XANES) and found that oxide coating is effective to prevent not only the mutual diffusion but also the degradation of the cathode induced by the formation of the oxygen defect. They selected two oxides, Li−Ge−O and Li−Zr−O, which are particularly effective for the mutual diffusion and the oxygen defect, respectively, and made a double coating layer on LiNi0.8Co0.15Al0.05O2 cathode so that the Li−Zr−O is inside. They found that the cell performance became better than the single coating. The relatively poor ionic conductivity of SE compared to liquid electrolyte is another challenge. Significant progress, however, has been made in the recent six years with the discovery of some sulfide-based compounds with higher ionic conductivities. R. Kanno’s group synthesized Li10GeP2S12 (LGPS) and demonstrated an ionic conductivity of 12 mS/ cm at room temperature.22 Substitution of Ge in LGPS by Sn or Si was suggested computationally to also give a high conductivity23 and recently confirmed experimentally.24−28 The Si-substituted version of LGPS was predicted to have the lowest activation energy of Li migration,23 and a recent Cl-modified version, Li9.54Si1.74P1.44S11.7Cl0.3, has demonstrated ionic conductivity as large as 25 mS/cm.29 Using hot pressing at 280 °C to reduce grain boundary resistance in Li7P3S11, Seino et al.9 improved the ionic conductivity from ∼2 to 17 mS/cm. These ionic conductivities make the solid conductors comparable to liquid electrolytes and, combined with lithium transport numbers near unity, make them extremely attractive for commercial sulfide based ASSBs. First-principles calculations have been used successfully to predict new SE materials and point at relevant structure attributes that can lead to high Li mobility.30−34 W. Richards et al.34 computationally predicted that Na10SnP2S12, obtained by substituting Na for Li in Li10SnP2S12, would be a good Na ion conductor. Na10SnP2S12 was subsequently synthesized and was shown to have an experimentally measured ionic conductivity and activation energy in good agreement with the prediction.34 (Recently, it was also suggested that this product is Na11Sn2PS12, not Na10SnP2S12.35) Y. Wang et al.36 proposed that an anion framework that is close to body-centered cubic (bcc), such as in LGPS and Li7P3S11, is optimal for Li+ conduction due to its interconnected network of energetically equivalent tetrahedral sites. Using this

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EXPERIMENTAL SECTION

To synthesize LiZnPS4, stoichiometric amounts of Li2S (99.9% Alfa Aesar), P2S5 (99% Sigma-Aldrich), and ZnS (99.99% Sigma-Aldrich) were mixed by using a planetary ball mill (Fritsch Pulverisette 7) at 400 rpm for 1000 min in a zirconia pot with zirconia balls. The mixture was sealed in an evacuated quartz tube and heated to 290 °C for 12 h in a muffle furnace followed by slow cooling to room temperature. For Li1+2xZn1−xPS4 synthesis, we chose two different sets of precursors. One was simply to mix designated ratios of Li2S, P2S5, and ZnS to form Li1+2xZn1−xPS4. The other was to use LiZnPS4 and amorphous Li3PS4 (a-LPS) by 1−x:x (mol), where a-LPS was synthesized by ball-milling of 75% Li2S and 25% P2S5 mixture at 400 rpm for 1000 min. All samples were sealed in evacuated quartz tubes and heated at 220−290 °C. The crystal structure and phase purity of as-synthesized Li1+2xZn1−xPS4 were evaluated by X-ray diffraction (PANalytical Empyrean) with Cu Kα radiation. Electronic states of P−S structure in the samples were measured by Raman spectroscopy (Jasco, RMP210) with a 532 nm YAG laser. For each measurement, samples were prepared in an Ar glovebox and sealed in a container with an X-ray or a laser window not to expose the samples to the ambient. Solid-state nuclear magnetic resonance (NMR) experiments were performed on Bruker Avance III consoles, at magnets with 14.1 T B0 2237

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Chemistry of Materials field strength corresponding to a 1H Larmor frequency of 600.13 MHz. 7Li magic angle spinning (MAS) NMR spectra and 31P MAS NMR were acquired ex situ at 233.2 and 242.9 MHz with a 2.5 mm MAS probe and at a 15 kHz spinning rate with π/2-(one-pulse) measurements with 2 and 10 s of a last-delay. All of the samples were prepared in an Ar-filled glovebox and prevented from contamination by oxygen and moisture. The ionic conductivity of Li1+2xZn1−xPS4 was obtained by electrochemical impedance spectroscopy (Autolab, FRA32M) between 1 MHz and 10 mHz. For this measurement, the Li1+2xZn1−xPS4 powder was pressed at 300 MPa, leading to a pellet of 1.3 cm diameter and ∼1−2 mm thick. The densities of the pellets were 1.7− 1.9 g/cm3 for all samples. The pellet was then sandwiched between two 0.05 mm thick indium foil electrodes and pressed at 80 MPa. We used indium foils for this purpose because indium is soft and can improve the attachment to the rough surfaces of a SE pellet by adding a high pressure. This assemblage was then sandwiched between two stainless steel disks, which were used as the current collectors, and placed in a test cell comprised of a stainless steel outer casing with a Teflon insulator.19 The cell was pressed at 22 MPa by using the screw to maintain the electrochemical contact during the measurement. The cyclic voltammetry (CV) was performed on an AUTOLAB PGSTAT M10. The potential was swept between −0.3 and +5 V at a scan rate of 10 mV/s. A 50 μm thick stainless steel plate was used as a working electrode, and a 0.1 mm thick lithium foil was used for the counter and reference electrode. The Li1+2xZn1−xPS4 (200 mg) was prepressed at 30 MPa, and the stainless steel plate and lithium foil were affixed on the pellet with a pressure of 300 MPa. And we put the sample pellet in the test cell above explained. The electrochemical performance of all-solid-state batteries with Li1+2xZn1−xPS4 as the SE was examined using custom pressurized test cells. The cathode was Li2O-ZrO2 coated LiNi0.8Mn0.05Co0.15O2 (LZONMC), and the anodes were either artificial graphite (spherical shape, D50 = ca. 15 μm), 0.2 mm thick indium foil (Nilaco), or Li4Ti5O12 (LTO) (D50 = ca. 0.4 μm). The coating of Li2O-ZrO2 on the surface of NMC (D50 = ca. 6 μm) was created with a sol−gel method developed by Ito et al.,19 where the LZO sol was prepared from 2-propanol, lithium methoxide, and zirconium(IV) tetrapropoxide (Zr(OC3H7)4) in the molar ratio 200:2:1. In this study, we used a rolling fluidized coating machine (SFD-01, Powrex) to spray the sol onto the surface of the NMC particles. After the coating, the sample was heated at 350 °C for 1 h under air. The cathode and anode powders were mixed with SE powder and carbon nanofiber (fiber diameter = ca. 200−300 nm, length = ca. 25 μm) in the weight ratio 60:35:5. To assemble the solid-state battery, 150 mg of the SE powder was prepressed at 30 MPa, and the cathode composite powder (20 mg) and graphite (12 mg) or LTO (30 mg) anode composite powder were uniformly spread on the 1.3 cm diameter pellet surface. In some tests, 0.2 mm thick indium foil was used as an anode. The assemblage was then pressed at 300 MPa and placed in the test cell in the way explained above. All test cells were assembled in an Ar-filled glovebox due to the sensitivity of the samples to moisture, and each test cell was sealed in a laminated bag with Ar gas, and taken out from the glovebox for the measurements.

Figure 1. XRD patterns for LiZnPS4, Li2Zn0.5PS4 made from Li2S, P2S5 and ZnS (a) and Li2Zn0.5PS4 from LiZnPS4 and a-Li3PS4 (b) and the simulated LiZnPS4 pattern obtained by VESTA.

Figure 2. Raman spectra for LiZnPS4, Li2Zn0.5PS4 (a) and (b).

former peak is attributed to [PS4]3−,39 evidencing that a large fraction of the phosphorus in LiZnPS4 forms tetrahedral [PS4]3−. The latter peak can be assigned to [P2S6]4−,39 which likely comes from a Li4P2S6 impurity, which cannot be detected in the XRD pattern but likely exists in amorphous form. For Li2Zn0.5PS4-(a), there are three major peaks in the spectrum at 380, 400, and 420 cm−1, which can be attributed to [P2S6]4−, [P2S7]4−, and [PS4]3−,39 respectively. This is consistent with the XRD observation, where a considerable amount of impurity was detected. The spectrum of Li2Zn0.5PS4-(b) was almost identical to that of LiZnPS4 without peaks for [P2S6]4− and [P2S7]4−. The actual position of the ∼420 cm−1 peaks for both Li2Zn0.5PS4-(a) and -(b) are shifted toward small wavenumber compared to LiZnPS4. It may be induced by the lattice expansion of Li2Zn0.5PS4, which will be shown later, and it made the wavelength of the symmetric stretching of P−S bonds longer. The results of XRD and Raman indicate that the Li2Zn0.5PS4 made from LiZnPS4 and a-LPS forms the I4̅ structure with little amount of impurity. There is another peak at around ∼250 cm−1 for LiZnPS4, and not in both Li2Zn0.5PS4-(a) and -(b). At this moment, we have

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RESULTS Figure 1 shows the XRD patterns of as-synthesized LiZnPS4 and two differently prepared Li2Zn0.5PS4 from (a) Li2S, P2S5, and ZnS and (b) LiZnPS4 and a-LPS. The overall structure of the three materials matches well with a structural model simulated with I4̅ symmetry.37 LiZnPS4 and Li2Zn0.5PS4-(b) are relatively phase-pure, showing negligible impurities. However, Li2Zn0.5PS4-(a) contains major impurity phases, indexed as Li4P2S6 and ZnS. Figure 2 shows the Raman spectra for LiZnPS4 and Li2Zn0.5PS4-(a) and -(b). A large peak at 422 cm−1 and a tiny peak at 385 cm−1 are seen in the spectrum of LiZnPS4. The 2238

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Chemistry of Materials not found the origin of the peak and cannot give an explanation as to why it does not exist in the nonstoichiometric samples. The reason why the Li2Zn0.5PS4-(a) contains much impurity while the Li2Zn0.5PS4-(b) does not is probably due to the relative inactivity of ZnS toward Li2S and P2S5, so that the formation of Li2Zn0.5PS4 does not go to completion easily. Raising the temperature to accelerate the reaction may not be possible, because LiZnPS4 decomposes into ZnS and Li4P2S6 at high temperature.37 By using LiZnPS4 and a-Li3PS4 as the starting materials, however, this difficulty can be circumvented. Figure 3 shows the XRD patterns of Li1+2xZn1−xPS4 (x = 0, 0.125, 0.25, 0.375, 0.5, 0.625, 0.75, 0.8, 0.9) synthesized from

Figure 4. Lattice parameters a/b and c for Li1+2xZn1−xPS4 obtained from the XRD pattern.

this component is unclear but possibly the resistance of the grain boundary or the resistance between Li2Zn0.5PS4 pellet and the In electrodes. The fitted curves are also shown in Figure 5a. They did not go through the origin because of the inductance component. We estimated σ25°C from Rbulk. σ25°C of Li2Zn0.5PS4-(b) (2.6 × 10−4 S/cm) was 6 times larger than that of Li2Zn0.5PS4-(a) (4.4 × 10−5 S/m), suggesting that a large amount of impurities lead to poor conductivity. The results of the fitting are listed in Table S2 in the Supporting Information. Figure 5b shows σ25°C and the activation energies (Ea) for various Li1+2xZn1−xPS4. Ea was calculated from the slope of the Arrhenius plot of σ(T) measured from T = 0 to 60 °C. Some typical Arrhenius plots of conductivities are shown in Figure 5c. For σ25°C and Ea, the errors of the fitting were ∼0.5 and ∼2%, respectively, for all samples. (Only the errors for Ea are drawn in Figure 5b, because the errors for σ25°C are too small to draw.) The results of the fitting of σ25°C at various x are also listed in Table S2 in the Supporting Information. The scattering of the values for the R2 component was very large, because, as is seen in Figure 5, this component did not clearly appear in the Cole− Cole plot. Li1+2xZn1−xPS4 shows negligible conductivity when x ≤ 0.2. At x = 0.375, σ25°C starts to increase accompanied by a decrease of Ea. This tendency is consistent with the XRD data, which shows the lattice parameter changing for x > 0.375. The highest σ25°C achieved is 5.7 × 10−4 S/cm at x = 0.625. When x ≥ 0.75, σ25°C decreases, which likely corresponds to low phase purity and poor crystallinity of the material. The σ25°C at x = 0.75 became smaller than that at x = 0.625, although Ea is smaller. This reason is not clear, but we presume that the number of movable Li+ became less as the purity and the crystallinity became worse over x ≥ 0.75. The XRD spectrum shows that the I4̅ structure was formed across most of the compositional range, but we cannot rule out the presence of a-LPS. To eliminate the possibility that the high conductivity of the synthesized materials is due to unreacted aLPS, which is also known to be a reasonably good ionic conductor on the order of 0.1 mS/cm,40−42 we performed NMR measurements on the samples.

Figure 3. XRD pattern for Li1+2xZn1−xPS4 (x = 0, 0.125, 0.25, 0.375, 0.5, 0.625, 0.75, 0.8, 0.9) and the simulated LiZnPS4 pattern obtained by VESTA.

LiZnPS4 and a-LPS. In some patterns, we see a broad peak at ∼20°. It is a background of the sample holder, which prevents the sample from exposing to the ambient, occasionally produced in our XRD apparatus. As x increases, the peak intensity decreases while the peak width broadens, indicating poor crystallinity. However, the material maintains the I4̅ structure, even at x = 0.8. The peak around 30° for x = 0.75 and 0.8 samples is indexed as unreacted LiZnPS4. At x = 0.9, it is not clear if the crystal structure is still the I4̅ structure. Using the I4̅ structure model, the lattice parameters for each of the Li1+2xZn1−xPS4 compositions were refined using Rietvelt analysis, and the results are shown in Figure 4. The actual values are listed in Table S1 in the Supporting Information. The a and c values are constant for x < 0.375 and gradually change as x increases (a increases and c decreases). Figure 5a shows the Nyquist plot of impedance spectra of two Li2Zn0.5PS4 made from Li2S, P2S5, and ZnS and LiZnPS4 and a-LPS (samples -(a) and -(b) as above) obtained at room temperature. Here we explain the analysis using the data for Li2Zn0.5PS4-(a). The conductivities (σ25°C) were extracted from the equivalent circuit (inset), where Rbulk is the bulk resistance of Li2Zn0.5PS4, appearing as the large semicircle between 1 MHz and 1 kHz in Figure 5a. R2 represents the small component appearing between 1 kHz and 10 Hz. The origin of 2239

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Figure 6. NMR spectra for 31P MAS NMR (a) and 7Li MAS NMR (b) of LiZnPS4, Li2Zn0.5PS4, Li2.5Zn0.25PS4, and a-LPS.

the ligand of PS4− is closer than in the x = 0 structure. Even though spectra for Li2Zn0.5PS4 and Li2.5Zn0.25PS4 samples contain several components, they cannot be characterized simply as mixtures of the crystalline LiZnPS4 and amorphous LPS phase, indicating that replacing the Zn atom with Li indeed formed a different phase. The peak at ∼90 ppm is probably unreacted LiZnPS4 consistent with the peak at ∼30° in the XRD pattern. Since we cannot see the XRD peak for the samples with x ≤ 0.625, we consider that these are almost single phase materials, as demonstrated by Li2Zn0.5PS4. Figure 6b shows the 7Li MAS NMR spectra for LiZnPS4, Li2Zn0.5PS4, Li2.5Zn0.25PS4, and a-LPS, respectively. There are three narrow peaks in the spectrum of LiZnPS4. In these peaks, the largest peak at 3 ppm is attributed to the LiZnPS4 phase, and the other two peaks at 2.2 and 1.4 ppm are probably impurity phases. There is only one peak at 1.75 ppm for a-LPS, evidencing that this a-LPS is uniform amorphous. The NMR spectra for Li2Zn0.5PS4 and Li2.5Zn0.25PS4 look similar, where a sharp peak appeared at around 2.15 ppm and a small mound at around 3 ppm. These features are probably due to Li1+2xZn1−xPS4 and unreacted residual LiZnPS4, respectively. Since there is no peak at 1.75 ppm, which is characteristic of aLPS, they do not contain a-LPS. From the NMR results, we confirm that the Li1+2xZn1−xPS4 samples synthesized in the present work are not a simple mixture of LiZnPS4 and a-LPS and that the samples with x ≤ 0.625 are almost single phase materials. The detailed interpretation of the NMR results, however, is rather difficult, and further experiment or modeling is necessary.

Figure 5. (a) The open circles and the open squares are Cole−Cole plots of the Li2Zn0.5PS4 (a) (open squares) and (b) (open circles), respectively. The solid lines represent the results of the fitting using the inset electrical equivalent circuit. (b) σ and Ea for Li1+2xZn1−xPS4 obtained from the fitting. (c) Some typical Arrhenius plots of conductivities.

Figure 6a shows the 31P MAS NMR spectra for LiZnPS4, Li2Zn0.5PS4, Li2.5Zn0.25PS4, and a-LPS, respectively. Increasing the x value (Zn composition decreases), the main peaks at 31P MAS NMR are upfield shifted from LiZnPS4. This suggests that 2240

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Chemistry of Materials Figure 7 shows the cyclic voltammogram for the Li1+2xZn1−xPS4 (x = 0.625) sample. There is almost no

Figure 8. Charge−discharge curves of the all-solid-state cells, whose anode is graphite (red), indium (blue), and LTO (black). The dashed line represents the charge−discharge curve of the cell whose SE is totally substituted with a-LPS. Figure 7. Cyclic voltammogram of Li2.25Zn0.375PS4.

potential of intercalation of Li into graphite anode, and it leads to overcharge of the cathode and no Li insertion into the anode. The cell with the indium anode shows ∼180 mAh/g in charge but 70 mAh/g in discharge. This indicates that Li2.25Zn0.375PS4 is also unstable against the indium anode, for which the alloying potential is 0.612 V (vs Li+/Li). When LTO is used, the discharge capacity obtained is 128 mAh/g−NMC. For comparison, we substituted the Li2.25Zn0.375PS4 in the cell to aLPS used in the synthesis of the Li1+2xZn1−xPS4, totally to form the configuration of LZO-NMC/a-LPS/LTO, and compared its performance. The voltage profile obtained in the cell is shown in Figure 8. The performance of the Li2.25Zn0.375PS4-based cell is almost identical to the a-LPS-based one. As a-LPS does not contain Zn, it can be more stable than Li2.25Zn0.375PS4 at low potentials. Given that they show identical charge and discharge voltage profiles, Li2.25Zn0.375PS4 must also be stable against LTO. Thus, Li2.25Zn0.375PS4 can be paired with a NMC cathode and LTO anode. In order to use the graphite anode, we also tested a cell in which a-LPS is used to inhibit the reaction between graphite and Li2.25Zn0.375PS4 similar to the approach taken by K. Takada et al.44 We mixed the graphite with a-LPS and the carbon nanofiber in the weight ratio 60:35:5 and used it as the anode. First, we pressed 75 mg of Li2.25Zn0.375PS4 powder at 40 MPa into a pellet. Next, we put 75 mg of a-LPS on the pellet and pressed at 40 MPa again to make a bilayer SE of Li2.25Zn0.375PS4 and a-LPS. Twenty mg of the Li2.25Zn0.375PS4-based cathode composite powder and 12 mg of the a-LPS-based graphite anode composite powder were uniformly spread on the Li2.25Zn0.375PS4 and a-LPS sides of the 1.3 cm bilayer pellet, respectively, and finally pressed to 400 MPa to form a graphite| a-LPS|Li2.25Zn0.375PS4|NMC battery. The charge−discharge curves of the graphite|a-LPS| i2.25Zn0.375PS4|NMC battery are shown in Figure 9. The test were carried out between 2.5 and 4.0 V at 50 μA/cm2 and 25 °C. The discharge capacity in the first cycle is about 144 mAh/ g−NMC (∼1.3 mAh/cm2), which is reasonable given the charge voltage (4.0 V). It shows good cyclability up to the 100th cycle. The capacity retention rate was 83% at the 100th cycle. There is a slope at the beginning of the first charge, probably because of the irreversible reaction on the graphite anode with a-LPS SE, which is also seen in ref 44. Although we obtained

oxidation reaction from OCV (∼1.5 V) to 5 V (vs Li+/Li), indicating that its electrochemical window at the cathode side reaches ∼5 V, which is anticipated to be feasible with typical 4 V class cathodes. During reduction, an appreciable amount of current begins flowing at ∼1 V (vs Li+/Li). We consider it is the onset of the reduction of Zn in Li1+2xZn1−xPS4 on the working electrode. Here we calculate the reduction potential of Zn in Li1+2xZn1−xPS4 using the grand potential phase diagram open to lithium in a similar method by Ong et al.43 In this case, the voltage is defined as the negative of the lithium chemical potential.43 The reduction of Zn into Zn−P alloys is expected to occur at 1.2 V (vs Li) with a corresponding uptake in lithium. The reaction is LiZnPS4 + 7Li → Li 2S + 0.25Zn3P2 + 0.25ZnP2

(1)

This is in excellent agreement with the results in Figure 7. Thus, the cathodic current in Figure 7 should be attributed to the reduction of Zn in Li1+2xZn1−xPS4 and subsequent Li−Zn alloying. In this case, the reduced Zn and Li−Zn alloy are conductive and do not work as a passivation layer. They move the reaction site deeply into the SE layer, and the reaction goes continuously. The reduction current drastically increases as the potential decreases, but almost no oxidation current appeared after it. This implies that the current may originate predominantly from the Zn reduction and/or Li−Zn alloying but not from Li deposition even below 0 V (vs Li+/Li). The above result suggests that Li1+2xZn1−xPS4 is electrochemically unstable against Li metal and other anodes with a redox potential