Mastering the interface for advanced all-solid-statelithium rechargeable batteriesYutao Lia,b,1, Weidong Zhoua,b,1, Xi Chena,b,1, Xujie Lüc, Zhiming Cuia,b, Sen Xina,b, Leigang Xuea,b, Quanxi Jiac,and John B. Goodenougha,b,2

aMaterials Science and Engineering Program, The University of Texas at Austin, Austin, TX 78712; bTexas Materials Institute, The University of Texas atAustin, Austin, TX 78712; and cCenter for Integrated Nanotechnologies, Los Alamos National Laboratory, Los Alamos, NM 87545

Contributed by John B. Goodenough, September 30, 2016 (sent for review August 24, 2016; reviewed by Ken Poeppelmeier and Jean-Marie Tarascon)

A solid electrolyte with a high Li-ion conductivity and a smallinterfacial resistance against a Li metal anode is a key componentin all-solid-state Li metal batteries, but there is no ceramic oxideelectrolyte available for this application except the thin-film Li-Poxynitride electrolyte; ceramic electrolytes are either easily re-duced by Li metal or penetrated by Li dendrites in a short time.Here, we introduce a solid electrolyte LiZr2(PO4)3 with rhombohe-dral structure at room temperature that has a bulk Li-ion conduc-tivity σLi = 2 × 10−4 S·cm−1 at 25 °C, a high electrochemical stabilityup to 5.5 V versus Li+/Li, and a small interfacial resistance for Li+

transfer. It reacts with a metallic lithium anode to form a Li+-con-ducting passivation layer (solid-electrolyte interphase) containingLi3P and Li8ZrO6 that is wet by the lithium anode and also wets theLiZr2(PO4)3 electrolyte. An all-solid-state Li/LiFePO4 cell with a polymercatholyte shows good cyclability and a long cycle life.

solid electrolyte | lithium anode | polymer catholyte | interfaces | NASICON

Arechargeable cell having a flammable organic liquid Li+

electrolyte has enabled the wireless revolution, but it is notable to power safely an electric road vehicle at a cost that iscompetitive with the gasoline-powered internal combustion en-gine (1–4). Safety concerns as well as cost, volumetric energydensity, and cycle life have prevented realization of a commer-cially viable electric road vehicle. To address this problem,considerable effort is being given to the development of a solidLi+ or Na+ electrolyte that is wet by a metallic lithium or sodiumanode and has an alkali ion conductivity σi > 10−4 S·cm−1 at thecell operating temperature Top, where a Top ≤ 25 °C is desired(5–10). Such a development would allow new as well as tradi-tional strategies for the cathode. Wetting of the solid electrolytesurface is desired not only because it prevents dendrite forma-tion and growth during plating of an alkali metal anode, but alsobecause wetting constrains the anode volume change in a charge/discharge cycle to be perpendicular to the anode/electrolyte in-terface, thereby allowing a long cycle life. Therefore, the shearmodulus of the electrolyte may not be critical where lithium wetsthe electrolyte surface.Ceramic oxide electrolytes offer a large energy gap between

their conduction and valence bands, which can allow realizationof a battery cell with a large energy separation between the an-ode and cathode chemical potentials without either reduction oroxidation of the electrolyte by an electrode (2). However, if analkali-metal anode reduces the solid electrolyte, formation of astable solid-electrolyte interphase (SEI) that conducts the workingLi+ or Na+ ion is acceptable if the Li+ or Na+ transfer across the SEIhas a low resistance. Although many ceramic solid Li+ electrolyteshave been investigated, they are easily reduced by Li metal and/orthey have failed to block dendrite formation and growth into theirgrain boundaries (SI Appendix, Fig. S1). However, the rhombo-hedral structure of the Na electrolyte Na1+3xZr2(SixP1–xO4)3 (11),NASICON, which was developed over 45 y ago, has recently beenused in a cell design in which a seawater cathode provides the so-dium of the anode (12).

The stability of the solid electrolyte on contact with a lithiumanode is a critical issue. If a lithium anode reduces the electro-lyte, (i) the electrolyte may become an electronic conductor,(ii) an interface layer may form that blocks Li+ transfer, or (iii) aninterface layer may form that conducts Li+ ions with a low im-pedance. The third situation forms with a LiZr2(PO4)3 electro-lyte. In this paper, we report that a Li+ electrolyte with theNASICON structure, LiZr2(PO4)3, can be fabricated by usingzirconium acetate as the precursor and spark plasma sintering(SPS); it forms a stable Li+-conducting SEI that is wet by ametallic lithium anode and also wets the electrolyte to provide asafe, all-solid-state Li/LiFePO4 cell operating at Top = 80 °C witha long cycle life; the LiFePO4 cathode particles are embedded ina polymer catholyte and carbon.

Results and DiscussionThe X-ray diffraction (XRD) results of as-prepared LiZr2(PO4)3are shown in Fig. 1A; the phase of LiZr2(PO4)3 depended on thestarting materials. Although different zirconium salts decomposedat the same temperature T < 400 °C (Fig. 1B), a pure LiZr2(PO4)3phase with rhombohedral structure could only be obtained withzirconium acetate; a triclinic phase with space group C-1 wasobtained with other zirconium materials. LiZr2(PO4)3 prepared bysolid-state reaction with ZrO2 as the starting material was repor-ted to change from triclinic to rhombohedral structure at 60 °C(13, 14). The rhombohedral LiZr2(PO4)3 phase with high Li-ionconductivity was stable at room temperature with zirconium ace-tate as the starting zirconium salt. The XRD refinement of therhombohedral LiZr2(PO4)3 phase is shown in Fig. 1C with re-liability factors (Rwp = 11.0%; RB = 2.3%). The rhombohedral

Significance

Realization of a safe, low-cost rechargeable lithium battery ofhigh energy density and long cycle life is needed for poweringan electric road vehicle and for storing electric power gener-ated by solar or wind energy. This urgent need has promptedefforts to develop a solid electrolyte with an alkali metal an-ode. Only now is it recognized that the key requirement iswetting of the electrolyte surface by the alkali-metal anode.We report a full rechargeable cell with a solid electrolyte that,although it is reduced by metallic lithium, forms a thin lithium–

electrolyte interface that is wet by the anode and wets theelectrolyte to give a small Li+ transfer resistance across theinterface.

Author contributions: Y.L. and J.B.G. designed research; Y.L., W.Z., X.C., Z.C., S.X., and L.X.performed research; Y.L., W.Z., X.L., and Q.J. analyzed data; and Y.L. and J.B.G. wrotethe paper.

Reviewers: K.P., Northwestern University; and J.-M.T., Collège de France.

The authors declare no conflict of interest.1Y.L., W.Z., and X.C. contributed equally to this work.2To whom correspondence should be addressed. Email: jgoodenough@mail.utexas.edu.

This article contains supporting information online at www.pnas.org/lookup/suppl/doi:10.1073/pnas.1615912113/-/DCSupplemental.

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phase has a space group R-3c with lattice parameters a = 8.8442 Åand c = 22.2645 Å, which is very close to those of LiZr2(PO4)3 at400 °C prepared by solid-sate reaction. A pure rhombohedralLiZr2(PO4)3 phase was also obtained by SPS. The electron-dis-persive spectroscopy mapping in SI Appendix, Fig. S2 indicates auniform distributions of Zr, P, and O elements in the LiZr2(PO4)3pellet. The LiZr2(PO4)3 pellets fired at 1,150 °C for 20 h in a boxfurnace and fired at 1,000 °C for 100 min by SPS have a density of85% (2.66 g·cm–3, SI Appendix, Fig. S3) and 99.9% (3.15 g·cm–3,Fig. 1D), respectively.A pure rhombohedral LiZr2(PO4)3 phase can be obtained

as a well-sintered ceramic at 1,200 °C by Y3+ or Ca2+ doping forZr4+; the rhombohedral structure remains unchanged from –70 °Cto 200 °C (15). The acetate route supplies particles with therhombohedral phase in a single firing at 900 °C; it is a simpler route,but the product undergoes a reversible change from rhombohedralto triclinic symmetry on cooling through 15 °C (SI Appendix, Fig.S3). Both products give comparable ionic conductivities at 80 °C.The impedance spectra of LiZr2(PO4)3 prepared with differ-

ent starting materials are shown in Fig. 2 and SI Appendix, Fig.S3. The rhombohedral LiZr2(PO4)3 pellets fired by conventionalsintering in a box furnace and by SPS at 1,000 °C for 10 min have,respectively, Li-ion conductivities of 2.2 × 10−5 S·cm–1 and 3.8 ×10−5 S·cm–1 at 25 °C, which is two to three orders of magnitudehigher than those of the room-temperature triclinic phases; the pelletfired by SPS method has a Li-ion conductivity of 1.8 × 10−4 S·cm−1

at 80 °C. The bulk Li-ion conductivity of rhombohedral LiZr2(PO4)3,as calculated from the distance between the zero point and the leftinterception of the semicircle with the Z′′ axis, was 2 × 10−4 S·cm−1

at 25 °C (SI Appendix, Fig. S3). The activation energy ofrhombohedral LiZr2(PO4)3 calculated from an Arrhenius plotover 300 K to 450 K in SI Appendix, Fig. S4 was 0.28 eV,

which is smaller than that (0.40 eV) of LiZr2(PO4)3 prepared bysolid-state reaction (15, 16). In rhombohedral LiZr2(PO4)3(Fig. 1C, Inset), 90% and 10% Li ions occupy, respectively, thesixfold disordered 36f M1 sites and the threefold disordered 18eM2 sites (13). The interstitial space is large enough for Li-iontransport, and higher Li-ion conductivity may be obtained bydoping with different valent ions to increase the Li-ion populationinside the framework.Whereas a garnet Li+ electrolyte can have a higher bulk σLi =

1 × 10−3 S·cm−1 at 25 °C (17), an insulating Li2CO3 surface layerforms on exposure to moist air; the interfaces of a symmetric Li/garnet/Li cell create a large resistance to Li+ transfer betweenthe anode and the electrolyte (about 1,700 Ω·cm−2, SI Appendix,Fig. S5) that dominates the total resistance of the electrolyte.In contrast, the SEI in the Li/LiZr2(PO4)3/Li interfaces with a400-μm-thick electrolyte pellet gave an interfacial resistance of650 Ω·cm−2 at 80 °C in the absence of an applied pressure on thecell. The much lower interfacial resistance of the Li/LiZr2(PO4)3/Licell makes LiZr2(PO4)3 a much better Li+ solid electrolytethan garnet.The electrochemical compatibility and stability of LiZr2(PO4)3

with Li metal was evaluated further at 80 °C with symmetric Li/LiZr2(PO4)3/Li cells by subjecting them to different currentdensities from 50 μA·cm–2 to 350 μA·cm–2. The symmetric cellshowed good cyclability at current densities less than 200 μA·cm–2,but the current densities above 200 μA·cm–2 increased the voltagea little after 50 h. At 50 μA·cm–2, the symmetric cell has a lowoverpotential of about 0.13 V and there is no evident voltage in-crease after 120 h; the interface between Li metal and LiZr2(PO4)3is very stable. Li/ LiZr2(PO4)3/Li cells were tested at 80 °C for upto 500 h. At 80 °C, a sealing problem with the coin cells resultedin oxidation of the lithium, but there was no evidence of failure

Fig. 1. (A) XRD patterns of LiZr2(PO4)3 prepared with different starting materials. (B) Thermogravimetric analysis (TGA) curves of LiZr2(PO4)3 with differentstarting materials. (C) Rietveld analysis of the XRD data of rhombohedral LiZr2(PO4)3. (D) SEM image of rhombohedral LiZr2(PO4)3 prepared by SPS at 1,000 °Cfor 10 min.

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by dendrite formation and growth. After cycling a Li/LiZr2(PO4)3/Licell for 10 h, a thin layer with black color formed only on thesurface of the LiZr2(PO4)3 pellet (SI Appendix, Fig. S6). Nonew diffraction peaks except those of the rhombohedral LiZr2(PO4)3 were observed with this thin layer (SI Appendix, Fig. S7),and no Li dendrites were found on the LiZr2(PO4)3 pellet aftercycling the symmetric cell at different current densities for300 h. The Raman inactive spectra of the black thin layer in Fig. 3Aindicated that the thin layer was amorphous. Fig. 2D shows thesurface of LiZr2(PO4)3 after cycling the Li/LiZr2(PO4)3/Li cell at50 μA·cm–2 for 100 h; the particle size was much smaller than that offresh LiZr2(PO4)3, but it still kept a dense surface structure. Therewere no Li metal dendrites on the Li metal surface after cycling thesymmetric cell at 150 μA·cm–2 for 20 h (Fig. 2E and SI Appendix,Fig. S8). The lithium metal wet well the surface of the thin surfacelayer on the LiZr2(PO4)3 electrolyte that formed on reaction witha Li electrode.To explore the composition of the thin interfacial layer that

formed during cycling at 80 °C, a Li metal foil on the surface of aLiZr2(PO4)3 pellet was heated from 25 °C to 350 °C (SI Ap-pendix, Fig. S6). For temperature below 300 °C, there was nochange of the LiZr2(PO4)3 pellet; a dense, black thin layer on thesurface of LiZr2(PO4)3 was formed when the pellet was heated at350 °C for 30 min. The XRD result of the pellet confirmed thedecomposition of LiZr2(PO4)3 to trigonal Li8ZrO6 and hexago-nal Li3P phases by the reaction

24Li+LiZr2ðPO4Þ3 → 2Li8ZrO6   +  3Li3P.

The XPS result of the black thin layer is shown in Fig. 3 C−F andSI Appendix, Fig. S9; the main peak at 284.8 eV in the C 1s

spectrum corresponds to adventitious carbon, and the small peakat 286.2 eV in SI Appendix, Fig. S9 may be from the residue ofthe polymer glue during the polishing process. No Li2CO3 peakat 288 eV was observed; LiZr2(PO4)3 is more stable againstmoist air than garnet Li7La3Zr2O12, which forms a Li-ion insulatingLi2CO3 layer on the particle surfaces during cooling in the firingprocess. The NASICON structure with strong P–O bonds usuallyhas high stability against moist air; for example, NASICONLi1.3Al0.3Ti1.7(PO4)3 is stable in air and in water (18, 19). ALi2CO3-related peak in the C 1s spectrum was from the reactionof Li metal with organic electrolyte in the glove box. The Zr4+ 3d3/2and 3d5/2 peaks at 185.5 eV and 183.13 eV in LiZr2(PO4)3shifted, respectively, to 184.8 eV and 181.7 eV after reaction,which is similar to the binding energy of Zr4+ in ZrO2. Afterreaction, one more peak at 127.13 eV in the P 2p spectrumcorresponds to P3– ions in Li3P, and the peak at 132.8 eV is fromthe P5+ ions in LiZr2(PO4)3. The O2– 2p peaks in LiZr2(PO4)3shifted from 531.25 eV to 530.3 eV after reaction, which is thesame as the O2– 2p binding energy in Li8ZrO6. No clear Li 1sbinding energy difference was observed in LiZr2(PO4)3 beforeand after reaction. Li3P in the thin layer is a good Li-ion con-ductor, and Li8ZrO6 with a layer structure may also have someLi-ion conductivity. After the reaction of LiZr2(PO4)3 with Limetal, a Li/LiZr2(PO4)3/Au cell was assembled with Au contact-ing the black thin layer; the cyclic voltammogam of the cell inSI Appendix, Fig. S10 showed that the black thin layer was

Fig. 2. (A) The impedance plots of LiZr2(PO4)3 at 25 °C and 80 °C. (B) The im-pedance plots Li/LiZr2(PO4)3/Li symmetric cell. (C) Cyclability of the Li/LiZr2(PO4)3/Lisymmetric cell for 300 h with different current densities. (D) Surface SEM imageof a LiZr2(PO4)3 pellet after cycling the cell for 100 h. (E) SEM image of Li metalafter charging the symmetric cell for 20 h at 150 μA·cm–2.

Fig. 3. (A) Raman spectra of the black SEI layer on the surface of LiZr2(PO4)3after cycling the Li/LiZr2(PO4)3/Li cell for 10 h. (B) XRD pattern of LiZr2(PO4)3after reaction with Li metal at 350 °C for 0.5 h. (C−F) XPS data of LiZr2(PO4)3before and after reaction with Li metal at 350 °C for 0.5 h.

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unstable at high voltage. The Li3P was reported to decomposeat voltage above 0.7 V, which resulted in a voltage increase inthe Li/LiZr2(PO4)3/Li cell at high current densities. For the appli-cation of LiZr2(PO4)3 in Li metal batteries, the thin layer onlyforms at the Li metal side, so the electrolyte will be stable duringthe charge/discharge process of an all-solid-state Li metal battery.The electrochemical stability window of LiZr2(PO4)3 is shown

in Fig. 4A; the two peaks near 0 V versus Li+/Li correspond to Limetal deposition and dissolution; there were no other redoxpeaks up to 5.5 V. An all-solid-state Li/LiZr2(PO4)3/LiFePO4battery was fabricated with the LiFePO4 cathode embedded in aLi-ion polymer electrolyte and carbon in a loading of 2 mg·cm–2.The polymer membrane with a melting point of 240 °C has aLi-ion conductivity of 1 × 10−4 S·cm−1 and an electrochemicalstability up to 4.7 V at 65 °C. The Li/LiZr2(PO4)3/LiFePO4battery had a small interfacial resistance with LiFePO4 of about300 Ω·cm−2 and a total resistance of 1,100 Ω·cm−2 at 80 °C(Fig. 4B), which is smaller than those of previously reported all-solid-state batteries (20). Fig. 4C shows the charge/dischargevoltage profiles at current densities of 50 μA·cm–2, 75 μA·cm–2,and 100 μA·cm–2 at 80 °C; the cells had a discharge capacity of140 mAh·g–1 and 120 mAh·g–1 with a cell polarization of 0.1 Vand 0.2 V at 50 μA·cm–2 and 100 μA·cm–2, respectively. A highcoulombic efficiency of 99.5 ± 0.5% over 40 cycles was obtained,which indicates that the Li/LiZr2(PO4)3 and LiZr2(PO4)3/LiFePO4interfaces were stable during cycling.

ConclusionsWe have prepared a stable rhombohedral NASICON LiZr2(PO4)3electrolyte at room temperature. A thin amorphous interfacial layercontaining Li8ZrO6 and Li3P formed on the LiZr2(PO4)3 surface byreaction with Li metal; this layer is wet by Li metal, which suppressesLi dendrite formation. LiZr2(PO4)3 has a Li-ion conductivity of 2 ·10−4 S·cm−1 at 80 °C, a small interfacial resistance against Li metaland a LiFePO4 cathode, and a large electrochemical window up to

5.5 V. An all-solid-state Li metal battery with LiZr2(PO4)3 as a solidelectrolyte contacting a Li metal anode showed no Li dendriteformation and good cycling performance with a Li insertion cathodeembedded in a polymer catholyte.

Materials and MethodsThe stoichiometric amounts of Li2CO3, (NH4)2HPO4, and different zirconiumsalts [Zr(AC)4, ZrOCl2, Zr(NO4)3 and ZrO2] were fired at 900 °C for 10 h, andthe obtained powders were ground and fired at 1,150 °C for 20 h in a Ptcrucible. The SPS pellets were obtained by firing the powders after 900 °C at1,000 °C for 10 min with a pressure of 50 Mpa. Powder XRD was used tomonitor the phase formation with a step size of 0.02°. A field emissionscanning electron microscope was used to obtain the fracture surface mi-crostructure of the pellet, and the distribution of elements was measured byenergy dispersive spectroscopy. Ionic conductivity was measured from 298 Kto 450 K with a Solarton Impedance Analyzer. Cyclic voltammetry was carriedout on an Auto Lab workstation at a scan rate of 0.5 mV·s−1. The Li/LiZr2(PO4)3/Lisymmetric cell was prepared by putting lithium foil on both sides of the LiZr2(PO4)3pellet, and the cell was cycled with different current densities in a Land in-strument. The LiFePO4 preparation process was the same as in our previousreport (20). To prepare the cathode of an all-solid-state LiFePO4/Li cell, theactive material LiFePO4 was mixed with carbon black, cross-linked poly-ethylene oxide, and LiTFSI (60:12:20:8 by weight) and ground in a mortar. Themixture was then dispersed in dimethylfluoride and stirred overnight. Theobtained slurry was spread evenly on a carbon-coated aluminum foil to pro-duce an electrode film with an active material loading of 2 mg·cm–2, whichwas dried at 90 °C for 12 h under vacuum. Then, 2,032 coin cells were fabri-cated in an argon-filled glove box with lithium foil as the anode. The cell wastested between 2.7 V and 3.8 V vs. Li+/Li in a Land instrument.

ACKNOWLEDGMENTS. This work was supported by National Science Foun-dation (NSF) Grant CBET-1438007 and the US Department of Energy, Officeof Basic Energy Sciences, Division of Materials Sciences and Engineering,under Award DESC0005397. The SPS processing at The University of Texas atAustin was conducted with an instrument acquired with the support of NSFAward DMR-1229131. The work at Los Alamos National Laboratory wasperformed, in part, at the Center for Integrated Technologies, an Office ofScience User Facility operated for the US Department of Energy Office ofScience.

Fig. 4. (A) A cyclic voltammogram of LiZr2(PO4)3 at a scanning rate of 0.5 mV·s–1. (B) The impedance plots of all-solid-state Li/LiZr2(PO4)3/LiFePO4 cell.(C) Charge and discharge voltage profiles and (D) cycling performance of an Li/LiZr2(PO4)3/LiFePO4 all-solid-state battery at 80 °C with different current densities.

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1. Tarascon JM, Armand M (2001) Issues and challenges facing rechargeable lithiumbatteries. Nature 414(6861):359–367.

2. Goodenough JB, Kim Y (2010) Challenges for rechargeable Li batteries. Chem Mater22(3):587–603.

3. Goodenough JB, Park KS (2013) The Li-ion rechargeable battery: A perspective. J AmChem Soc 135(4):1167–1176.

4. Ellis BL, Lee KT, Nazar LF (2010) Positive electrode materials for Li-Ion and Li-batteries.Chem Mater 22(3):691–714.

5. Murugan R, Thangadurai V, Weppner W (2007) Fast lithium ion conduction in garnet-type Li7La3Zr2O12. Angew Chem Int Ed Engl 46(41):7778–7781.

6. Adachi GY, Imanaka N (1996) Aono H fast Li+ conducting ceramic electrolytes. AdvMater 8(2):127–135.

7. Catti M (2007) First-principles modeling of lithium ordering in the LLTO (LixLa2/3-x/3TiO3)superionic conductor. Chem Mater 19(16):3963–3972.

8. Thangadurai V, Narayanan S, Pinzaru D (2014) Garnet-type solid-state fast Li ionconductors for Li batteries: Critical review. Chem Soc Rev 43(13):4714–4727.

9. Lü X, et al. (2014) Li-rich anti-perovskite Li3OCl films with enhanced ionic conductivity.Chem Commun (Camb) 50(78):11520–11522.

10. Li Y, et al. (2016) Fluorine-doped antiperovskite electrolyte for all-solid-state lithium-ion batteries. Angew Chem Int Ed Engl 55(34):9965–9968.

11. Goodenough JB, Hong HYP, Kafalas JA (1976) Fast Na+-ion transport in skeletonstructures. Mater Res Bull 11(2):203–220.

12. Kim Y, Kim H, Park S, Seo I, Kim Y (2016) Na ion conducting ceramic as solid elec-

trolyte for rechargeable seawater batteries. Electrochim Acta 191:1–7.13. Catti M, Comotti A, Di Blas S (2003) High-temperature lithiummobility in α-LiZr2(PO4)3

NASICON by neutron diffraction. Chem Mater 15(8):1628–1632.14. Catti M, Morgante N, Ibberson RM (2000) Order–disorder and mobility of Li+ in the

β′-and β-LiZr2(PO4)3 ionic conductors: A neutron diffraction study. J Solid State Chem

152(2):340–347.15. Li Y-T, Liu M-J, Liu K, Wang C-A (2013) High Li+ conduction in NASICON-type

Li1+xYxZr2−x(PO4)3 at room temperature. J Power Sources 240:50–53.16. Xie H, Goodenough JB, Li Y-T (2011) Li1.2Zr1.9Ca0.1(PO4)3, a room-temperature Li-ion

solid electrolyte. J Power Sources 196(18):7760–7762.17. Li Y-T, Han J-T, Wang C-A, Xie H, Goodenough JB (2012) Optimizing Li+ conductivity

in a garnet framework. J Mater Chem 22(30):15357–15361.18. Hasegawa S, et al. (2009) Study on lithium/air secondary batteries—Stability of

NASICON-type lithium ion conducting glass–ceramics with water. J Power Sources

189(1):371–377.19. Shimonishi Y, et al. (2011) A study on lithium/air secondary batteries—Stability of the

NASICON-type lithium ion conducting solid electrolyte in alkaline aqueous solutions.

J Power Sources 196(11):5128–5132.20. Zhou W, et al. (2016) Plating a dendrite-free lithium anode with a polymer/ceramic/

polymer sandwich electrolyte. J Am Chem Soc 138(30):9385–9388.

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