Structural and Electronic-State Changes of a Sulfide Solid Electrolyte

May 3, 2017 - Department of Interdisciplinary Environment, Graduate School of Human and Environmental Studies, Kyoto University, Yoshida-nihonmatsu-ch...
3 downloads 20 Views 2MB Size
Download PDF
Article pubs.acs.org/cm

Structural and Electronic-State Changes of a Sulfide Solid Electrolyte during the Li Deinsertion−Insertion Processes Takashi Hakari,† Minako Deguchi,† Kei Mitsuhara,‡ Toshiaki Ohta,‡ Kohei Saito,§ Yuki Orikasa,§ Yoshiharu Uchimoto,§ Yoshiyuki Kowada,∥ Akitoshi Hayashi,*,† and Masahiro Tatsumisago†,* †

Department of Applied Chemistry, Graduate School of Engineering, Osaka Prefecture University, 1-1 Gakuen-cho, Naka-ku, Sakai, Osaka 599-8531, Japan ‡ SR Center, Ritsumeikan University, 1-1-1 Noji-Higashi, Kusatsu, Shiga 525-8577, Japan § Department of Interdisciplinary Environment, Graduate School of Human and Environmental Studies, Kyoto University, Yoshida-nihonmatsu-cho, Sakyo-ku, Kyoto 606-8501, Japan ∥ Department of Natural Sciences, Hyogo University of Teacher Education, 942-1 Shimokume, Kato city, Hyogo 673-1494, Japan ABSTRACT: All-solid-state batteries with sulfide solid electrolytes (SEs) are promising next-generation energy storage devices that are safe and have a cycle life and energy densities. To improve their electrochemical performances, we investigated the electrochemical reactions of the SEs and their structural changes. The detailed changes in the structure and electronic states have not yet been measured experimentally because of the difficulty in observing the microscopic area formed at the interface between the electrode materials and the SEs. Thus, we prepared composite electrodes composed of Li3PS4 SE and carbon to increase the electrochemical reaction area. The structural and electronic-state changes in Li3PS4 during the Li deinsertion−insertion processes were revealed using X-ray diffraction and Raman, X-ray photoelectron, and X-ray absorption spectroscopies. We found that the sulfide ions in Li3PS4 contribute to charge compensation during the charge−discharge processes. The S−S bonds between PS4 units associate and dissociate via the retention of the covalent bonds between the P atoms and S−S bonds. Although the anion redox behavior is generally discussed with regard to transition metal chalcogenides, here we first report the reversible association and dissociation of S−S bonds in typical element compounds. This knowledge contributes to an improved understanding of the anion redox reactions and the interfacial resistances between high-voltage positive electrodes and SEs.

■

cm−1 at room temperature. However, the oxide SEs require cosintering to increase the level of contact between the active materials and the SE, which leads to side reactions causing high interfacial resistance. Sulfur-based SEs such as Li10GeP2S12 (LGPS),5 Li7P3S11 glass-ceramic,23,24 and Li9.54Si1.74P1.44S11.7Cl0.325 are preferable because of their high ionic conductivities that are comparable to those of liquid electrolytes and better mechanical properties that improve the physical contact with the electrodes.7,26 To utilize active materials such as high-voltage oxide active materials and Li metal, the electrochemical window of the SE is one of the most important features. SEs have extremely wide electrochemical windows of 0−5 V versus Li+/Li, as measured by cyclic voltammetry (CV) measurements.5,9,23 However, there are many inconsistent reports concerning the electrochemical windows determined by using CV. On the low-voltage

INTRODUCTION

Lithium-ion batteries (LIBs) are used in energy storage applications.1 However, the applications of large LIBs for energy storage and transport applications, such as in aircraft and automobiles, are limited by the safety concerns regarding the flammable organic liquid electrolytes used in LIBs.2,3 Allsolid-state lithium batteries, which have nonflammable inorganic solid electrolytes (SEs), have attracted attention because of their excellent safety.4−12 In addition to the improved safety, all-solid-state lithium batteries are particularly promising with regard to energy density; consequently, allsolid-state lithium batteries could be used as high-voltage positive electrodes and high-capacity sulfur and lithium electrodes.13−17 Furthermore, bulk-type all-solid-state batteries, in which the composite electrodes are composed of a mixture of active material particles, SEs, and conductive additives, are compact, increasing the packing efficiency of the cells. Recently, the development of SEs with high ionic conductivities has accelerated rapidly. NASICON-type18,19 and garnet-type oxide20−22 SEs have ionic conductivities of around 10−3 S © 2017 American Chemical Society

Received: February 9, 2017 Revised: May 2, 2017 Published: May 3, 2017 4768

DOI: 10.1021/acs.chemmater.7b00551 Chem. Mater. 2017, 29, 4768−4774

Article

Chemistry of Materials side, Whiteley et al. investigated27 the interfacial resistance between LGPS and Li metal by attaching Li foils to the LGPS pellets. The resistance of the cell increased with storage time, suggesting that a decomposition layer grew between LGPS and Li metal. Oh et al.17 reported the decomposition product at the interface between LGPS and Li metal, and the formation of Li2S was assessed using X-ray diffraction (XRD) after electrochemical cycling. Recently, Wenzel et al. have investigated the interfacial reactions between Li metal and SEs such as LGPS and Li7P3S11 with in situ X-ray photoelectron spectroscopy (XPS).28,29 At the interfaces, decomposition layers composed of Li compounds such as Li2S and Li3P are formed, suggesting that the SEs are reduced. On the highvoltage side, the layer with a high interfacial resistance that forms between the high-voltage oxide active material and the SE is one of the reasons for the low power densities of all-solidstate batteries. Various oxide coatings on the active materials have been investigated to decrease the interfacial resistance.30−33 The interfacial resistance is considered to be caused by the formation of a space charge layer and chemical side reaction at the interface.4,34−36 Recently, the oxidation of SEs has been proposed as another reason for the high interfacial resistance between the high-voltage oxide active materials and the SEs.37−41 The electrochemical windows of SEs calculated using density functional theory (DFT) are extremely narrow compared to those measured by CV.37−39 In addition, the oxidation of SE (the extraction of Li from the SEs) at the interface between the high-voltage electrodes and SEs has been predicted by density functional molecular dynamics (DF-MD) simulations.40 Therefore, in CV measurements, redox decomposed layers are thought to be formed at and limited to the interface between the SE and the current collector because the redox currents measured by CV are directly proportional to the total amount of decomposition layer.41,42 In fact, minute redox currents are observed in the magnified CV curves. 43 Furthermore, thin decomposition layers (several tens of nanometers thick) have been observed at the interfaces between the SE and Li by in situ XPS.28,29 To clearly observe the redox currents, Han et al. prepared SE−carbon composite electrodes by ball-milling, which resulted in a large contact area between the SEs and the electronic conductive material; this increased the size of the decomposition layers.41 In the CV measurements, the all-solid-state cell with the prepared composite electrodes exhibits noticeable oxidation currents, corresponding to the calculated electrochemical windows. However, there are some studies concerning the involvement of SEs in the electrochemical reaction and their use as an active material in all-solid-state cells.42,44−48 Shin et al.44 found that the TiS2−Li3PS4 composite electrodes prepared by ball-milling had extra capacity, and they suggested that the Li−Ti−P−S amorphous phase functions as an active material in the composite electrode. Lin et al.45 proposed a series of lithiumconducting sulfur-rich compounds, Li3PS4+n, and applied them to all-solid-state Li/S cells. Furthermore, Nagata et al. prepared46 S−P2S5-based positive electrodes without Li+ ions, and the all-solid-state cells containing the electrodes exhibited a high capacity and rate performance. Ramgasamy et al.47 reported additional capacity beyond the theoretical limit in all-solid-state Li−CFx batteries with composite electrodes composed of CFx, carbon, and β-Li3PS4. They concluded that the additional capacity originates from the electrochemical discharge of β-Li3PS4. We have prepared Li2S active material−

Li3PS4 glass-carbon composite electrodes by ball-milling and determined that the Li3PS4 in the composite functions not only as an ionic conduction path to Li2S but also as an active material.42 Finally, Han et al. demonstrated48 a proof-ofconcept single-material battery utilizing LGPS−carbon composites in both electrodes. Although interesting phenomena about SEs have been observed, the reaction mechanism of SEs has not been elucidated in detail. To further improve the interface modification between the high-voltage active materials and the SEs and the utilization of the SEs as an active material, it is necessary to understand their Li deinsertion−insertion mechanism. In previous studies, the structural changes have been mainly predicted using DFT and DF-MD simulations35,37−40 because of the empirical difficulty in observing the microscopic area formed at the interface. Therefore, the structural and electronic-state changes have not yet been experimentally measured. In this study, we investigated the structural and electronicstate changes to sulfide SEs during Li deinsertion and insertion by using XRD, Raman spectroscopy, XPS, and X-ray absorption fine structure (XANES). To clearly observe the redox layer, we increased the size of the total redox layer by preparing SE− carbon composite electrodes with a large contact area between the SE and carbon. The Li3PS4 glass, which is composed of only PS4 tetrahedral units,49 was utilized as a model SE to gain a simple understanding of the structural and electronic-state changes.

■

EXPERIMENTAL SECTION

Preparation of the Li3PS4−Carbon Composite. To prepare the Li3PS4 glass powders, a mixture of reagent-grade Li2S (Mitsuwa Chemicals Co., Ltd.) and P2S5 (Aldrich, 99%) powders was placed into a ZrO2 pot (45 mL) with 500 ZrO2 balls (diameter of 4 mm) and then mechanically milled with a planetary ball mill (Fritsch Pulverisette 7) for 10 h at 510 rpm. The Li3PS4−carbon composite electrode with a weight ratio of 70:30 was obtained by ball-milling of a mixture of Li3PS4 and acetylene black (AB, Denki Kagaku Kogyo) at 510 rpm for 1 h in a ZrO2 pot (45 mL) with 160 ZrO2 balls (diameter of 4 mm) in a dry Ar-filled glovebox. Fabrication of All-Solid-State Cells with the Li3PS4−AB Composite Electrodes. The all-solid-state cells of In/Li3PS4/ Li3PS4−AB composite electrodes were fabricated as follows. The Li3PS4−AB composite electrode was directly used as a working electrode, and the Li3PS4 glass was used as a SE layer. The obtained composite electrode (4−5 mg) and the SE layer (80 mg) were set in a polycarbonate tube (diameter of 10 mm) and pressed together under a pressure of 360 MPa. An indium foil with a thickness of 100 μm was then placed on the surface of the SE side of the bilayer pellet as a counter electrode. The three-layer pellet was then sandwiched between two stainless steel disks, which acted as current collectors. Finally, the electrochemical solid-state cell was obtained by pressing at 120 MPa. An electrochemical test of the cells was conducted at a constant current density of 0.13 mA cm−2 (30 mA g−1) at 25 °C in an Ar atmosphere using a charge−discharge measuring device (BTS2004, Nagano Co.). Characterization of the Li3PS4−AB Composite Electrodes. To analyze the structural changes of Li3PS4 in the working electrode, we used XRD and Raman spectroscopy measurements of the working electrode before and after the initial charge−discharge process using an X-ray diffractometer (Cu Kα, Ultima IV; Rigaku Corp.) and a laser Raman spectrometer (He−Ne laser, LabRAM HR-800; Horiba Ltd.). The dispersion state of AB in the composite electrodes was investigated with a cross-sectional field emission scanning electron microscope (FE-SEM, SU8220; Hitachi Co.). Cross sections of the working electrode were thinned down for analysis using an ion milling system (IM4000; Hitachi Co.). The electronic states of sulfur and phosphorus in the Li3PS4−AB composite electrode before and after 4769

DOI: 10.1021/acs.chemmater.7b00551 Chem. Mater. 2017, 29, 4768−4774

Article

Chemistry of Materials the charge−discharge tests were analyzed by XPS (K-Alpha; Thermo Fisher Scientific Co.) with a monochromatic Al Kα source (1486.6 eV). The observed binding energies were calibrated with the adventitious C 1s peak to 284.7 eV. The samples were mounted on a sample stage in a dry Ar-filled glovebox and transferred to an analysis chamber using an Ar-filled transfer vessel. XANES spectra at S and P K-edge were measured at BL-13 of the Synchrotron Radiation Center of Ritsumeikan University. The spectra were recorded in total electron yield mode, wherein the sample current produced by the excitation of electrons of different energies was measured. The samples for S and P K-edge XANES measurements were also mounted on a sample stage in a dry Ar-filled glovebox and transferred to the chamber using an Arfilled transfer vessel.

spherical, and the particles are 30−50 nm in size, indicating that the intrinsic shape of AB is maintained even after ball-milling. In addition, the AB particles are dispersed uniformly in the Li3PS4 matrix, suggesting a large contact area between Li3PS4 and AB. Figure 1c shows the S2p and P2p XPS spectra of Li3PS4 glass and Li3PS4−AB composite. The S2p spectrum of Li3PS4 glass is composed of doublet peaks (S2p3/2 and S2p1/2). The peak energy separation (1.2 eV) in each doublet is set during peak fitting. The doublet peaks can be attributed to a nonbridging sulfur (S−).52,53 In the spectrum of the composite, other doublet peaks slightly exist at the higher-binding energy side in addition to the peaks of S−. In the P2p spectrum of Li3PS4 glass, doublet peaks that can be attributed to P2p3/2 and P2p1/2 are observed. The main peak at 131.5 eV (P2p3/2) can be attributed to P in the PS43− units. Although some small peaks are also present at lower and higher binding energies, there are no apparent differences in the P2p spectra of the Li3PS4 glass and the Li3PS4−AB composite. Electrochemical Characterization. The first charge− discharge curve of the all-solid-state cell of the In/Li3PS4 glass/Li3PS4−AB composite is shown in Figure 2. The obtained

■

RESULTS AND DISCUSSION Characterization of the Li3PS4−AB Composite Electrodes. The Raman spectra, cross-sectional FE-SEM image, and XPS spectra of the Li3PS4−AB composite electrode are displayed in Figure 1. Figure 1a shows the Raman spectra of

Figure 2. Charge−discharge curve of an all-solid-state cell with the Li3PS4−carbon composite electrode at 25 °C and 0.13 mA cm−2 (30 mA g−1). The arrows show the XANES measurement points.

capacity was normalized by a weight of Li3PS4 glass in the composite. The cell containing the Li3PS4−AB composite shows an initial charge capacity of ∼250 mAh g−1 and a discharge capacity of ∼150 mAh g−1, corresponding to 1.5 mol of deintercalated Li+ ions, where 1.0 mol of Li+ ions is intercalated in the Li3PS4 active material in the initial charge− discharge processes. Electrode samples for XRD, Raman spectroscopy, and XPS were electrochemically prepared in the fully charged and discharged states. Electrode samples for the XANES were also prepared at different stages (indicated in Figure 2) of charge−discharge capacity. XRD Patterns and Raman Spectra in the Initial Charge−Discharge Processes. Figure 3 shows the (a) XRD patterns and (b) Raman spectra of the Li3PS4−AB composite electrode before and after the initial charge− discharge processes. In the XRD patterns, Si powder was added as an internal standard. A halo pattern was observed, and new crystalline peaks were not observed, suggesting that Li3PS4 glass maintains an amorphous structure even after charging and discharging. In all Raman spectra, Raman bands attributed to PS43− in the Li3PS4 glass are present. The Li3PS4 glass after the first cycle has a local structure similar to that before the cycle. These structural analyses indicate that the Li3PS4 glass is not crystalline and its local structure is maintained after the initial charge−discharge process.

Figure 1. (a) Raman spectra of carbon, Li3PS4 glass, and the Li3PS4− carbon composite. (b) Cross-sectional FE-SEM image of the Li3PS4− carbon composite electrode. (c) P 2p and S 2p XPS spectra of Li3PS4 and the Li3PS4−carbon composite electrode.

AB (A), Li3PS4 (B), and the Li3PS4−AB composite (C). For reference, spectrum A shows the characteristic features of carbon with the D band at 1348 cm−1 and the G band at 1590 cm−1.50 Spectrum B is that of Li3PS4 glass. The main peak at 420 cm−1 can be attributed to the PS43− structural unit of the Li3PS4 glass,51 and this peak is unchanged in the spectra of the composite (C), indicating that the local bonding environment in the Li3PS4 glass has not been changed by the ball-milling process. Figure 1b shows a back-scattered electron (BSE) image of the cross section of the Li3PS4−AB composite pellet. Black spots are visible in a lighter gray matrix. The BSE image shows that the contrast corresponds to a difference in the atomic number. This allows us to observe the AB distribution from the difference in the contrast between Li3PS4 and AB in the composite. Thus, the lighter gray and black parts can be attributed to Li3PS4 and AB, respectively. AB particles are 4770

DOI: 10.1021/acs.chemmater.7b00551 Chem. Mater. 2017, 29, 4768−4774

Article

Chemistry of Materials

at 162.7 eV is assigned to the bridging sulfur, and the association and dissociation of the S−S bonds occur during the charging and discharging processes. XANES Spectra in the Initial Charge−Discharge Processes. S K-edge XANES spectra of the Li3PS4−AB electrode before and after the charge−discharge processes are shown in panels a and b of Figure 5. The peak intensity of the

Figure 3. (a) XRD patterns and (b) Raman spectra of the Li3PS4− carbon composite electrode before and after the charge−discharge processes.

XPS Spectra in the Initial Charge−Discharge Processes. S2p and P2p XPS spectra of the Li3PS4−AB composite electrodes before and after the initial charge−discharge processes are shown in Figure 4. In the S2p XPS spectra, the

Figure 5. S K-edge XANES spectra of the Li3PS4−carbon composite electrode during (a) charging and (b) discharging. P K-edge XANES spectra of the Li3PS4−carbon composite electrode during (c) charging and (d) discharging.

spectra at approximstely 2471−2472 eV increased and decreased during the charge and discharge processes, respectively, suggesting that the changes to the electronic states are reversible. The intense peak at approximately 2471− 2472 eV can be attributed to a transition from the S 1s orbital to the antibonding S 3pσ* orbital of dumbbell-shaped S22−, which have been reported in the FeS2 crystal structure.54 The increase in peak intensity indicates that the number of S−S bonds increased during the charging process. This result agrees with that of the S 2p XPS spectrum shown in Figure 4. The P K-edge XANES spectra of the Li3PS4−AB electrode before and after the charge−discharge process were also analyzed to confirm the electronic-state changes of the phosphorus atoms, as shown in panels c and d of Figure 5. In terms of the charging process, the intensity of the peak at 2148 eV decreases with increasing peak intensities in the S K-edge XANES spectrum. These observations are very interesting, and we will now discuss the reaction mechanism. First, we compared the changes in the P K-edge XANES spectra with the oxidation of Li2MnO3 in terms of the electronic structure because the top valence band of Li2MnO3 is formed of oxygen orbitals; thus, electrons are removed from the oxygen orbitals during the charging process. Oishi et al. have reported the changes to the electronic structure of Li2 MnO 3 during charging and discharging by using XANES measurements.55 During the charging process, the reduction of Mn and the formation of peroxide and superoxide are characterized by the Mn L-edge and O K-edge XANES spectra. This suggests that electrons of

Figure 4. P 2p and S 2p XPS spectra of the Li3PS4−carbon composite electrode before and after charge−discharge processes. The S 2p XPS spectrum of Na2S4 is shown as a reference.

intensities of the doublet peaks at higher binding energies (red) are increased and decreased during the charging and discharging, respectively. On the other hand, the P 2p XPS spectra do not change significantly during the charge−discharge processes. These results suggest that sulfide ions contribute to charge compensation during the charge−discharge processes. In the density of states (DOS) of Li3PS4 glass, S 3p electrons in Li3PS4 glass are occupied at the top of the valence band,49 indicating that sulfur redox reactions mainly occur in the Li3PS4 glass. The electrochemical redox species considered from the XPS results agree with that from the DOS. The S2p spectrum of Na2S4 was measured as a reference sample. Crystalline Na2S4 contains two sulfur species: terminal nonbridging S− and bridging sulfur S0 in S−S bonds combined with the S− in Na2S4. The peak at 162.7 eV (S2p3/2) of Li3PS4 is close to the S2p3/2 (162.8 eV) peak of the bridging sulfur of Na2S4. The S2p3/2 peak 4771

DOI: 10.1021/acs.chemmater.7b00551 Chem. Mater. 2017, 29, 4768−4774

Article

Chemistry of Materials

oxide positive electrode materials might be developed by increasing the size of the contact area between typical element compounds and electronic conductive materials. Figure 6 shows schematic illustrations of the Li3PS4−AB composite electrode (a) and the experimentally observed

antibonding 2pπ* of peroxide transfer to the Mn 3d orbitals. In our case, a peak that can be attributed to the transition from the 1s orbital to the 3pπ* orbital formed upon dissociation of P and S−S bonds should be observed in the S K-edge spectra after the charging process. However, the peak at lower binding energies, around 2468 eV,56 that can be attributed to disulfur radical S2•− ion was not observed in the S K-edge spectra upon charging (Figure 5a). Therefore, the decrease in the peak intensity of P K-edge spectra is not explained by the reduction of P caused by the dissociation of S22−. This interpretation is supported by the P2p XPS spectrum after charging because the P2p spectrum does not change significantly as shown in Figure 4. However, currently, an origin of the changes in the P K-edge XANES spectra is not known because of a lack of theoretical data about the P K-edge XANES. Pascal et al. simulated the S K-edge spectra of polysulfide using first-principles molecular dynamics.57 They observed an unexpected decrease in the peak intensity attributed to the bridging sulfur in Li2S8, which was assigned to delocalized electrons in the core-excited state along the sulfur chain. Thus, the decrease in the peak intensity for the P K-edge spectra during charging might be attributable to the delocalization of electrons in the core excited states along the P−S−S−P chains. The electrochemical window of Li3PS4 calculated by using DFT indicates that there are two possible oxidation reactions of Li3PS4.37−39 One is the conversion of Li3PS4 to S and P2S5,37,38 and the other is the formation of S and Li4P2S6 from Li3PS4.39 In these oxidation reactions, the peak at 164 eV (S2p3/2)58 that can be attributed to sulfur should be observed in the XPS spectra after the charging process. However, the peak is not observed, and thus, these reactions can be ignored from the results of S2p XPS spectra. According to the results of the spectroscopy data presented above, the most likely structure is considered to result from the formation of S−S bonds between the PS4 units. This scenario is supported by the interfacial structure of LiFePO4/Li3PS4 predicted by DF-MD simulations.40 Furthermore, covalent bonds between P and the S−S bond are maintained even after Li+ ions are extracted from the Li3PS4. Therefore, the changes to the structure and electronic states of the SE as the active material are reversible. In fact, the reversible redox reactions of transition metal dichalcogenide are promoted by the degree of metal−ligand (cation−anion) covalent bonding (e.g., FeS2, CoS2, and NiS2).59,60 In addition, in oxide electrode materials such as Li2RuO3, strong covalent bonds between Ru and peroxide allow the reversible anion redox reactions.61,62 Reversible redox reactions of dichalcogenides have generally been reported and discussed in transition metal compounds, and our data are the first to demonstrate the reversible association and dissociation of S−S bonds of typical element compounds. The covalent bonds between P and S−S bonds are maintained because the transfer of charge from the 3pπ* orbital of the S−S bond to P is energetically unfavorable because of the formation of unstable P4+ ions. However, this speculation requires further verification by investigating the bonding and electronic states during the charge−discharge processes. This phenomenon could also be applied to the oxide anion redox reaction. In fact, recently, a new class of amorphous active material of LixNiyPOz has been shown to have a capacity of 350 mAh g−1 after 200 cycles in all-solid-state thin film batteries.63 The electron numbers of the electrochemical reaction calculated from the reversible capacities suggest the occurrence of anion redox reactions, although the authors did not discuss anion redox processes. Typical element

Figure 6. Schematic illustrations of the (a) Li3PS4−AB composite electrode and (b) experimentally observed structure of the redox layer at the interface between Li3PS4 and AB in the all-solid-state In/Li3PS4/ Li3PS4−AB cell during the charging process. The black sphere, red ring around the black sphere, and yellow background indicate Li3PS4 (the solid electrolyte), AB, and the redox layer, respectively.

structure of the redox layer (b) at the interface between Li3PS4 and AB in the In/Li3PS4/Li3PS4−AB cell during the charging process; the structure is proposed on the basis of the charge− discharge measurements, XRD, Raman, FE-SEM, XPS, and XANES analyses. The yellow background, black sphere, and red part around the black sphere show the Li3PS4 solid electrolyte, AB, and the redox layer, respectively. Li3PS4 is maintained after ball-milling with AB, as shown by the Raman and XPS spectra in Figure 1. In addition, the all-solid-state cell with the Li3PS4− AB composite electrodes had the charge and discharge capacities shown in Figure 2. Li3PS4 itself is oxidized and then reduced, and the size of the total redox layer of Li3PS4 is increased by increasing the interface area between Li3PS4 and AB, as shown in Figure 6a. The results of XRD, Raman, XPS, and S K-edge XANES investigations of the samples after the charging process suggest that S−S bonds are formed between the PS4 units, as shown in Figure 6b. Furthermore, the P and S K-edge XANES spectra indicate that the covalent bonds between P and S−S bonds are maintained even after Li+ ions are extracted from the Li3PS4; therefore, reversible changes in the structure and electronic state of SE as the active material occurred. At the interface between the high-voltage positive electrode and SEs, the Li extracted layer would have a high resistance because of the low Li content in the layer. To inhibit the extraction of Li+ ions, it is necessary to increase the Li chemical potential by using a lithium-doped oxide coating on the high-voltage positive electrodes.

■

CONCLUSIONS We characterized the redox layer and the structural and electronic-state changes of Li3PS4 during Li deinsertion− insertion processes using XRD, Raman, XPS, and XANES. The charge compensation of sulfide ions and the association and dissociation of the S−S bonds between the PS4 units occurred during the Li deinsertion−insertion processes. In addition, the reversible structural and electronic-state changes of Li3PS4 as an active material were allowed by the strong covalent bonds between P and the S−S bonds even though Li+ ions were extracted from the SEs. This results in an improved understanding of the anion redox chemistry and the interfacial 4772

DOI: 10.1021/acs.chemmater.7b00551 Chem. Mater. 2017, 29, 4768−4774

Article

Chemistry of Materials

(14) Unemoto, A.; Yasaku, S.; Nogami, G.; Tazawa, M.; Taniguchi, M.; Matsuo, M.; Ikeshoji, Y.; Orimo, S. Development of Bulk−Type All−Solid−State Lithium−Sulfur Battery using LiBH4 electrolyte. Appl. Phys. Lett. 2014, 105, 083901. (15) Yamada, T.; Ito, S.; Omoda, R.; Watanabe, T.; Aihara, Y.; Agostini, M.; Ulissi, U.; Hassoun, J.; Scrosati, B. All − Solid−State Lithium−Sulfur Battery Using a Glass−Type P2S5−Li2S Electrolyte: Benefits on Anode Kinetics. J. Electrochem. Soc. 2015, 162, A646− A651. (16) Yubuchi, S.; Ito, Y.; Matsuyama, T.; Hayashi, A.; Tatsumisago, M. 5 V class LiNi0.5Mn1.5O4 Positive Electrode Coated with Li3PO4 Thin Film for All−Solid−State Batteries using Sulfide Solid Electrolyte. Solid State Ionics 2016, 285, 79−82. (17) Oh, G.; Hirayama, M.; Kwon, O.; Suzuki, K.; Kanno, R. BulkType All Solid-State Batteries with 5 V Class LiNi0.5Mn1.5O4 Cathode and Li10GeP2S12 Solid Electrolyte. Chem. Mater. 2016, 28, 2634−2640. (18) Fergus, J. W. Ion Transport in Sodium Ion Conducting Solid Electrolytes. Solid State Ionics 2012, 227, 102−112. (19) Aono, H. Ionic Conductivity of Solid Electrolytes Based on Lithium Titanium Phosphate. J. Electrochem. Soc. 1990, 137, 1023− 1027. (20) Murugan, R.; Thangadurai, V.; Weppner, W. Fast Lithium Ion Conduction in Garnet-Type Li7La3Zr2O12. Angew. Chem., Int. Ed. 2007, 46, 7778−7781. (21) Ohta, S.; Kobayashi, T.; Asaoka, T. High Lithium Ionic Conductivity in the Garnet−Type Oxide Li7−XLa3(Zr2−X, NbX)O12 (X = 0−2). J. Power Sources 2011, 196, 3342−3345. (22) Allen, J.; Wolfenstine, J.; Rangasamy, E.; Sakamoto, J. Effect of Substitution (Ta, Al, Ga) on the Conductivity of Li7La3Zr2O12. J. Power Sources 2012, 206, 315−319. (23) Mizuno, F.; Hayashi, A.; Tadanaga, K.; Tatsumisago, M. New, Highly Ion−Conductive Crystals Precipitated from Li2S-P2S5 Glasses. Adv. Mater. 2005, 17, 918−921. (24) Seino, Y.; Ota, T.; Takada, K.; Hayashi, A.; Tatsumisago, M. A Sulphide Lithium Super Ion Conductor is Superior to Liquid Ion Conductors for Use in Rechargeable Batteries. Energy Environ. Sci. 2014, 7, 627−631. (25) Kato, Y.; Hori, S.; Saito, T.; Suzuki, K.; Hirayama, M.; Mitsui, A.; Yonemura, M.; Iba, H.; Kanno, R. High-Power All−Solid−State Batteries using Sulfide Superionic Conductors. Nat. Energy 2016, 1, 16030. (26) Sakuda, A.; Hayashi, A.; Tatsumisago, M. Sulfide Solid Electrolyte with Favorable Mechanical Property for All-Solid-State Lithium Battery. Sci. Rep. 2013, 3, 2261. (27) Whiteley, J. M.; Woo, J. H.; Hu, E.; Nam, K.-W.; Lee, S.-H. Empowering the Lithium Metal Battery through a Silicon−Based Superionic Conductor. J. Electrochem. Soc. 2014, 161, A1812−A1817. (28) Wenzel, S.; Weber, D. A.; Leichtweiss, T.; Busche, M. R.; Sann, J.; Janek, J. Interphase Formation and Degradation of Charge Transfer Kinetics between a Lithium Metal Anode Highly Crystalline Li7P3S11 Solid Electrolyte. Solid State Ionics 2016, 286, 24−33. (29) Wenzel, S.; Randau, S.; Leichtweiss, T.; Weber, D. A.; Sann, J.; Zeier, W. G.; Janek, J. Direct Observation of the Interfacial Instability of the Fast Ionic Conductor Li10GeP2S12 at the Lithium Metal Anode. Chem. Mater. 2016, 28, 2400−2407. (30) Woo, J. H.; Trevey, J. E.; Cavanagh, A. S.; Choi, Y. S.; Kim, S. C.; George, S. M.; Oh, K. H.; Lee, S.-H. Nanoscale Interface Modification of LiCoO2 by Al2O3 Atomic Layer Deposition for Solid− State Li Batteries. J. Electrochem. Soc. 2012, 159, A1120−A1124. (31) Okada, K.; Machida, N.; Naito, M.; Shigematsu, T.; Ito, S.; Fujiki, S.; Nakano, M.; Aihara, Y. Preparation and Electrochemical Properties of LiAlO2−Coated Li(Ni1/3Mn1/3Co1/3)O2 for All-SolidState Batteries. Solid State Ionics 2014, 255, 120−127. (32) Yada, C.; Ohmori, A.; Ide, K.; Yamasaki, H.; Kato, T.; Saito, T.; Sagane, F.; Iriyama, Y. Dielectric Modification of 5V−Class Cathodes for High−Voltage All−Solid−State Lithium Batteries. Adv. Energy Mater. 2014, 4, 1301416. (33) Kim, J.; Kim, O.; Park, C.; Lee, G.; Shin, D. Electrochemical Properties of Li1+xCoO2 Synthesized for All−Solid−State Lithium Ion

resistance between the high-voltage positive electrodes and the SEs.

■

AUTHOR INFORMATION

Corresponding Authors

*E-mail: (A.H.) [email protected]. *E-mail: (M.T.) [email protected]. ORCID

Akitoshi Hayashi: 0000-0001-9503-5561 Notes

The authors declare no competing financial interest.

■

ACKNOWLEDGMENTS This research was supported by the Japan Science and Technology Agency (JST), Advanced Low Carbon Technology Research and Development Program (ALCA), Specially Promoted Research for Innovative Next-Generation Batteries (SPRING) project. T.H. is grateful for a Grant-in-Aid and a fellowship for the Japan Society for the Promotion of Science (JSPS).

■

REFERENCES

(1) Goodenough, J. B.; Kim, Y. Challenges for Rechargeable Li Batteries. Chem. Mater. 2010, 22, 587−603. (2) Ellis, B. L.; Lee, K. T.; Nazar, L. F. Positive Electrode Materials for Li-Ion and Li−Batteries. Chem. Mater. 2010, 22, 691−714. (3) Manthiram, A. Materials Challenges and Opportunities of Lithium Ion Batteries. J. Phys. Chem. Lett. 2011, 2, 176−184. (4) Ohta, N.; Takada, K.; Zhang, L.; Ma, R.; Osada, M.; Sasaki, T. Enhancement of the High−Rate Capability of Solid−State Lithium Batteries by Nanoscale Interfacial. Adv. Mater. 2006, 18, 2226−2229. (5) Kamaya, N.; Homma, K.; Yamakawa, Y.; Hirayama, M.; Kanno, R.; Yonemura, M.; Kamiyama, T.; Kato, Y.; Hama, S.; Kawamoto, K.; Mitsui, A. A Lithium Superionic Conductor. Nat. Mater. 2011, 10, 682−686. (6) Aboulaich, A.; Bouchet, R.; Delaizir, G.; Seznec, V.; Tortet, L.; Morcrette, M.; Rozier, P.; Tarascon, J. M.; Viallet, V.; Dollé, M. A New Approach to Develop Safe All−Inorganic Monolithic Li−Ion Batteries. Adv. Energy Mater. 2011, 1, 179−183. (7) Hayashi, A.; Noi, K.; Sakuda, A.; Tatsumisago, M. Superionic Glass−Ceramic Electrolytes for Room−Temperature Rechargeable Sodium Batteries. Nat. Commun. 2012, 3, 856. (8) Yersak, T. A.; Macpherson, H. A.; Kim, S. C.; Le, V. D.; Kang, C. S.; Son, S. B.; Kim, Y. H.; Trevey, J. E.; Oh, K. H.; Stoldt, C.; Lee, S.H. Solid State Enabled Reversible Four Electron Storage. Adv. Energy. Mater. 2013, 3, 120−127. (9) Liu, Z.; Fu, W.; Payzant, E. A.; Yu, X.; Wu, Z.; Dudney, N. J.; Kiggans, J.; Hong, K.; Rondinone, A. J.; Liang, C. Anomalous High Ionic Conductivity of Nanoporous β−Li3PS4. J. Am. Chem. Soc. 2013, 135, 975−978. (10) Wang, Y.; Richards, W. D.; Ong, S. P.; Miara, L. J.; Kim, J. C.; Mo, Y.; Ceder, G. Design Principles for Solid−State Lithium Superionic Conductors. Nat. Mater. 2015, 14, 1026−1031. (11) Park, K. H.; Oh, D. Y.; Choi, Y. E.; Nam, Y. J.; Han, L.; Kim, J.Y.; Xin, H.; Lin, F.; Oh, S. M.; Jung, Y. S. Solution−Processable Glass LiI−Li4SnS4 Superionic Conductors for All−Solid−State Li−Ion Batteries. Adv. Mater. 2016, 28, 1874−1883. (12) Li, Y.; Zhou, W.; Xin, S.; Li, S.; Zhu, J.; Lu, X.; Cui, Z.; Jia, Q.; Zhou, J.; Zhao, Y.; Goodenough, J. B. Fluorine−Doped Antiperovskite Electrolyte for All−Solid−State Lithium−Ion Batteries. Angew. Chem., Int. Ed. 2016, 55, 9965−9968. (13) Nagao, M.; Hayashi, A.; Tatsumisago, M. Fabrication of Favorable Interface between Sulfide Solid Electrolyte and Li Metal Electrode for Bulk−Type Solid−State Li/S Battery. Electrochem. Commun. 2012, 22, 177−180. 4773

DOI: 10.1021/acs.chemmater.7b00551 Chem. Mater. 2017, 29, 4768−4774

Article

Chemistry of Materials Batteries with Li2S−P2S5 Glass−Ceramics Electrolyte. J. Electrochem. Soc. 2015, 162, A1041−A1045. (34) Sakuda, A.; Hayashi, A.; Tatsumisago, M. Interfacial Observation between LiCoO2 Electrode and Li2S−P2S5 Solid Electrolytes of All−Solid−State Lithium Secondary Batteries Using Transmission Electron Microscopy. Chem. Mater. 2010, 22, 949−956. (35) Haruyama, J.; Sodeyama, K.; Han, L.; Takada, K.; Tateyama, Y. Space−Charge Layer Effect at Interface between Oxide Cathode and Sulfide Electrolyte in All−Solid−State Lithium−Ion Battery. Chem. Mater. 2014, 26, 4248−4255. (36) Yokokawa, H. Thermodynamic Stability of Sulfide Electrolyte/ Oxide Electrode Interface in Solid−State Lithium Batteries. Solid State Ionics 2016, 285, 126−135. (37) Zhu, Y.; He, X.; Mo, Y. Origin of Outstanding Stability in the Lithium Solid Electrolyte Materials: Insight from Thermodynamic Analyses Based on First−principles Calculations. ACS Appl. Mater. Interfaces 2015, 7, 23685−23693. (38) Zhu, Y.; He, X.; Mo, Y. First Principles Study on Electrochemical and Chemical Stability of Solid Electrolyte-Electrode Interfaces in All−Solid−State Li−ion Batteries. J. Mater. Chem. A 2016, 4, 3253−3266. (39) Richards, W. D.; Miara, L. J.; Wang, Y.; Kim, J. C.; Ceder, G. Interface Stability in Solid−State Batteries. Chem. Mater. 2016, 28, 266−273. (40) Sumita, M.; Tanaka, Y.; Ikeda, M.; Ohno, T. Charged and Discharged States of Cathode/Sulfide Electrolyte Interfaces in All− Solid−State Lithium Ion Batteries. J. Phys. Chem. C 2016, 120, 13332− 13339. (41) Han, F.; Zhu, Y.; He, X.; Mo, Y.; Wang, C. Electrochemical Stability of Li10GeP2S12 and Li7La3Zr2O12 Solid Electrolytes. Adv. Energy. Mater. 2016, 6, 1501590. (42) Hakari, T.; Nagao, M.; Hayashi, A.; Tatsumisago, M. All− Solid−State Lithium Batteries with Li3PS4 Glass as Active Material. J. Power Sources 2015, 293, 721−725. (43) Hayashi, A.; Muramatsu, H.; Ohtomo, T.; Hama, S.; Tatsumisago, M. Improvement of Chemical Stability of Li3PS4 Glass Electrolytes by Adding MxOy (M = Fe, Zn, and Bi) Nanoparticles. J. Mater. Chem. A 2013, 1, 6320−6326. (44) Shin, B. R.; Nam, Y. J.; Kim, J. W.; Lee, Y.-G.; Jung, Y. S. Interfacial Architecture for Extra Li+ Storage in All-Solid-State Lithium Batteries. Sci. Rep. 2015, 4, 5572. (45) Lin, Z.; Liu, Z.; Fu, W.; Dudney, N. J.; Liang, C. Lithium Polysulfidophosphates: A Family of Lithium−Conducting Sulfur−Rich Compounds for Lithium-Sulfur Batteries. Angew. Chem., Int. Ed. 2013, 52, 7460−7463. (46) Nagata, H.; Chikusa, Y. Transformation of P2S5 into a Solid Electrolyte with Ionic Conductivity at the Positive Composite Electrode of All−Solid−State Lithium−Sulfur Batteries. Energy. Technol. 2014, 2, 753−756. (47) Rangasamy, E.; Li, J.; Sahu, G.; Dudney, N.; Liang, C. Pushing the Theoretical Limit of Li−CFx Batteries: A Tale of Bifunctional Electrolyte. J. Am. Chem. Soc. 2014, 136, 6874−6877. (48) Han, F.; Gao, T.; Zhu, Y.; Gaskell, K. J.; Wang, C. A. Battery Made from a Single Material. Adv. Mater. 2015, 27, 3473−3483. (49) Ohara, K.; Mitsui, A.; Mori, M.; Onodera, Y.; Shiotani, S.; Koyama, Y.; Orikasa, Y.; Murakami, M.; Shimoda, K.; Mori, K.; Fukunaga, T.; Arai, H.; Uchimoto, Y.; Ogumi, Z. Structural and Electronic Features of Binary Li2S−P2S5 Glasses. Sci. Rep. 2016, 6, 21302. (50) Ferrari, A. C.; Robertson, J. Resonant Raman Spectroscopy of Disordered, Amorphous, and Diamondlike Carbon. Phys. Rev. B: Condens. Matter Mater. Phys. 2001, 64, 075414. (51) Mizuno, F.; Hayashi, A.; Tadanaga, K.; Tatsumisago, M. High Lithium Ion Conducting Glass−Ceramics in The System Li2S−P2S5. Solid State Ionics 2006, 177, 2721−2725. (52) Hayashi, A.; Tatsumisago, M.; Minami, T.; Miura, Y. Structural Investigation of 95(0.6Li2S·0.4SiS2)·5Li4SiO4 Oxysulfide Glass by Using X−Ray Photoelectron Spectroscopy. J. Am. Ceram. Soc. 1998, 81, 1305−1309.

(53) Wang, Y.; Matsuyama, T.; Deguchi, M.; Hayashi, A.; Nakao, A.; Tatsumisago, M. X-Ray Photoelectron Spectroscopy for Sulfide Glass Electrolytes in The Systems Li2S−P2S5 and Li2S−P2S5−LiBr. J. Ceram. Soc. Jpn. 2016, 124, 597. (54) Womes, M.; Karnatak, R. C.; Esteva, J. M.; Lefebvre, I.; Allan, G.; Olivier-Fourcades, J. O.; Jumas, J. C. Electronic Structure of FeS and FeS2: X-Ray Absorption Spectroscopy and Band Structure Calculation. J. Phys. Chem. Solids 1997, 58, 345−352. (55) Oishi, M.; Yamanaka, K.; Watanabe, I.; Shimoda, K.; Matsunaga, T.; Arai, H.; Ukyo, Y.; Uchimoto, Y.; Ogumi, Z.; Ohta, T. Direct Observation of Reversible Oxygen Anion Redox Reaction in Li−Rich Manganese Oxide, Li2MnO3, Studied by Soft X−Ray Absorption Spectroscopy. J. Mater. Chem. A 2016, 4, 9293−9302. (56) Pascal, T. A.; Pemmaraju, C. D.; Prendergast, D. X−Ray Spectroscopy as a Probe for Lithium Polysulfides Radicals. Phys. Chem. Chem. Phys. 2015, 17, 7743−7753. (57) Pascal, T. A.; Wujcik, K. H.; Velasco-Velez, J.; Wu, C.; Teran, A. A.; Kapilashrami, M.; Cabana, J.; Guo, J.; Salmeron, M.; Balsara, N.; Prendergast, D. X-Ray Absorption Spectra of Dissolved Polysulfides in Lithium−Sulfur Batteries from First−Principles J. Phys. J. Phys. Chem. Lett. 2014, 5, 1547−1551. (58) Matsuyama, T.; Deguchi, M.; Mitsuhara, K.; Ohta, T.; Mori, T.; Orikasa, Y.; Uchimoto, Y.; Kowada, Y.; Hayashi, A.; Tatsumisago, M. Structure Analyses using X-Ray Photoelectron Spectroscopy and XRay Absorption Near Edge Structure for Amorphous MS3 (M:Ti, Mo) Electrodes in All-Solid-State Lithium Batteries. J. Power Sources 2016, 313, 104−111. (59) Jobic, S.; Brec, R.; Rouxel, J. Occurrence and Characterization of Anionic Bondings in Transition Metal Dichalcogenides. J. Alloys Compd. 1992, 178, 253−283. (60) Rouxel, J. Anion−Cation Redox Competition and the Formation of New Compounds in Highly Covalent Systems. Chem. - Eur. J. 1996, 2, 1053−1059. (61) Sathiya, M.; Rousse, G.; Ramesha, K.; Laisa, C. P.; Vezin, H.; Sougrati, M. T.; Doublet, M.-L.; Foix, D.; Gonbeau, D.; Walker, W.; Prakash, A. S.; Ben Hassine, M.; Dupont, L.; Tarascon, J.-M. Reversible Anionic Redox Chemistry in High−Capacity Layered− Oxide Electrodes. Nat. Mater. 2013, 12, 827−835. (62) Saubanere, M.; McCalla, E.; Tarascon, J.-M.; Doublet, M.-L. The Intriguing Question of Anionic Redox in High−Energy Density Cathodes for Li-Ion Batteries. Energy Environ. Sci. 2016, 9, 984−991. (63) Sabi, Y.; Sato, S.; Hayashi, S.; Furuya, T.; Kusanagi, S. A New Class of Amorphous Cathode Active Material LixMyPOz (M= Mi, Cu, Co, Mn, Au, Ag, Pd). J. Power Sources 2014, 258, 54−60.

4774

DOI: 10.1021/acs.chemmater.7b00551 Chem. Mater. 2017, 29, 4768−4774