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Recent Advancements in Li-ion Conductors for All-Solid-State Li-ion Batteries Yedukondalu Meesala, Anirudha Jena, Ho Chang, and Ru-Shi Liu ACS Energy Lett., Just Accepted Manuscript • DOI: 10.1021/acsenergylett.7b00849 • Publication Date (Web): 30 Oct 2017 Downloaded from http://pubs.acs.org on October 30, 2017

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ACS Energy Letters

Recent Advancements in Li-ion Conductors for All-Solid-State Li-ion Batteries Yedukondalu Meesala,† Anirudha Jena,‡ Ho Chang, ‡,* and Ru-Shi Liu†,‡,* †

Department of Chemistry, National Taiwan University, Taipei 106, Taiwan.

‡

Department of Mechanical Engineering and Graduate Institute of Manufacturing Technology,

National Taipei University of Technology, Taipei 106, Taiwan.

*Corresponding authors E-mails: [email protected] (RSL); Fax: +886-2-33668671; Tel: +886-2-33661169 [email protected] (HC); Fax: +886-2-27712171; Tel: +886-2-27317191

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ABSTRACT: Inorganic solid lithium ion conductors are potential candidates as replacement for conventional organic electrolytes for safety concerns. However, achieving a Li-ion conductivity comparable to that in existing liquid electrolytes (>1m S cm−1) remains a challenge in solid-state electrolytes. One of the approaches to achieve a desirable conductivity is doping of various elements into the lattice framework. Our discussion on the structure and conductivity of crystalline Li-ion conductors includes description on NASICON (NAtrium Super Ionic CONductor)-type conductors, garnettype conductors; perovskite-type conductors, and LISICON (Lithium Super Ionic CONductor)-type conductors. Moreover, we discuss various strategies currently used to enhance ionic conductivity including theoretical approaches; ultimately optimizing electrolyte/electrode interface and improving cell performance.

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The continuous depletion of fossil fuels, increasing oil prices, and the need to mitigate CO2 emissions have stimulated intensive research on alternative energy technologies based on renewable and clean sources. Among the renewable energy sources, solar and wind energies are the prominent energy options; however, the intensity of sunlight and wind constantly changes with locations. Thus, alternative resources must be stored. Rechargeable Li-ion batteries (LIBs) are one of the promising candidates for storage of alternative energy resources. Moreover, the LIBs have received tremendous attention due to their significant role in today’s consumer electronics market, such as mobile electronic devices, laptops, digital cameras, and electric vehicles.1 Within the battery class, LIBs, have outperformed Pb-acid, Ni-Cd, and Ni-MH batteries in terms of specific energy and power as represented in the Ragone plot (Figure 1).2 However, commercial LIBs contain toxic and flammable organic liquid as electrolyte, resulting in serious safety issues. Moreover, growth of Li-dendrite within the electrode materials due to such organic electrolytes lead to short-circuit and finally death of LIBs. Additionally, the limited electrochemical window of organic liquid electrolyte limits the choice of electrodes in LIBs.3-5 To address these problems, replacement of organic liquid with inorganic solid electrolytes (SEs) will improve safety and will ensure powering a wide range of electronic devices from portable to heavy locomotives. Also, use of SEs in place of organic electrolytes will suppress formation and growth of Li-dendrite during battery cycles.6 The desirable criteria for SEs for their implementation in real application is the ionic conductivity, which should be exceeding 10−3 S cm−1 at room temperature. Thermal, chemical, and mechanical stability will ensure the compatibility of the SEs with electrode materials and this will improve the SE-electrode interface stability and hence


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Figure 1. Ragone plot for various rechargeable batteries showing superiority of the Libased battery systems as the desired energy storing devices.

minimize the interfacial resistance. Besides, for the optimal ionic conductivity of SEs, the size of the mobile ion should fit in to the size of the migration channel (bottleneck) and the bottleneck should be smooth without very wide or narrow sections.7 The structural parameters can be finely tuned by dopants and substitutions with ions with different ionic radii into the lattice structure, for the suitable size of the bottleneck in the Li-ion conduction pathway (Figure 2a). The dopant ions with larger size would cause lattice expansion, thus increase the bottleneck size, and reduce the activation energy and improve the conductivity. Among the most studied fast ion conductors, NASICON (NAtrium Super Ionic CONductor)-type, garnet-type, perovskite-type, and LISICON (Lithium Super Ionic CONductor)-type inorganic ceramic SEs have displayed ionic conductivities in the order


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a)

b)

Figure 2. (a) Total ionic conductivities of selected solid-state lithium-ion conductors at room temperature. (b) Ionic conductivities of selected solid-state lithium-ion conductors with variation of temperature.

of 10−3 S cm−1 at room temperature (Figure 2a).8-10 Furthermore, these conductors exhibit outstanding stability within the electrochemical window of 0–9 V vs. Li/Li+ as revealed by cyclic voltammetry measurements.11 Figure 2b shows the Arrhenius plots of the ionic conductivities of the selected inorganic SEs.9 Knauth et al.12 published a review on the structures and conductivities for each class of SE. Bachman et al.13 published a review on the ion-transport mechanisms and fundamental properties of solid-state electrolytes. NASICON-type Li-ion conductors The general molecular formula of NASICON-type conductors is AxBy(PO4)3, where A and B are monovalent and multivalent cations, respectively. The crystallographic structure of the NASICON NaA2IV(PO4)3 (A = Ge, Ti, Sn, Hf, Zr) was first reported in 1968.14 The term NASICON has been coined by Hong and Goodenough for the framework Na1+xZr2P3-xSixO12 (0 ≤ x ≤ 3) in 1971.15-16 These Na super-ionic conductors


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display thermal and chemical stability and high ion diffusion properties. NASICON materials form three-dimensional network structures, with [ByP3O12]−x covalent skeletons built via corner sharing of BO6 octahedra and PO4 tetrahedra along the c-axis direction. NASICONs generally crystallize in thermally stable rhombohedral structures with R3c space group, although triclinic phases have been reported in compositions NaA2(PO4)3 with A = Sn, Zr, Hf.17 In rhombohedral phases, mobile A cations in the skeletons are distributed in two types of interstitial sites (Mʹ and M″) for charge compensation. M′ sites are surrounded by six oxygen atoms (octahedral symmetry) and situated between two adjacent [ByP3O12]−x units along the c- axis, whereas the M″ sites are surrounded by 10 oxygen atoms and symmetrically distributed around the three-fold axis of the structure. The M″ sites are located between two columns of [ByP3O12]−x units along the c-axis. The mobile A cations migrate through bottlenecks from one site to another, and the size of the bottlenecks depends on the nature of the skeleton ions and concentration of mobile ions at interstitial sites. Consequently, the ionic conductivity of NASICON materials depends on the concentration of Li-ions and composition of the NASICON framework. Among the members of the NASICON family LiM2(PO4)3 (M = Zr, Hf, Sn, Ti and Ge), LiZr2(PO4)3 (LZP), LiTi2(PO4)3 (LTP), and LiGe2(PO4)3 (LGP) are the most promising materials being investigated in recent years.18-19 LiZr2(PO4)3 displays a complex polymorphism, and the phase of the LZP system depends on the synthetic parameters. LiZr2(PO4)3 synthesized at high temperature (>1100 °C) adopts a rhombohedral α phase (R3c) at room temperature and undergoes a phase transition to triclinic α' phase at low temperature ( 6 V vs. Li/Li+ than the Nb-contained garnet Li5La3Nb2O12.63 The Liion conductivity can be enhanced significantly by partial substitution of La3+ or M5+ sites with aliovalent cations. The Ba-substituted garnet Li6BaLa2Ta2O12 displayed a high ionic conductivity of 4 × 10−5 S cm−1 at room temperature and an activation energy of 0.4 e V. The garnet system shows that the bulk and total conductivities are nearly of the same magnitude, indicating a low grain boundary resistance.63 The conductivity was further enhanced upon substitution with Y or In. The In-substituted garnet Li5.5La3Nb1.75In0.25O12 prepared at 950 °C displayed an enhanced bulk ionic conductivity of 1.8 × 10−4 S cm−1 at 50 °C with a low activation energy of 0.51 eV.64 The solid-state garnet doped with yttrium and lithium (Li6.5La3Nb1.25Y0.75O12) displayed the highest bulk conductivity of 2.7 × 10−4 S cm−1 at 25 °C.65 The yttrium-doped samples were chemically stable up to 400 °C–600 °C after water treatment, as well as chemically compatible in the presence of the high-voltage cathode materials Li2CoMn3O8 and Li2FeMn3O8.65 Li7La3Zr2O12 (LLZO) with garnet-related structure has been widely studied in recent years because of its good ionic conductivity at room temperature and stability against lithium metal. In 2007, Murugan et al.66 first reported the synthesis of the garnet-type LLZO through


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Figure 11. Crystal structure of tetragonal Li7La3Zr2O12 with connectivity pattern of Td (8a and 16e) and Oh (16f and 32g) cages projected on 2D.

substitution of Zr4+ for M5+ in Li5La3M2O12; this LLZO showed an ionic conductivity of 2.44 × 10−4 S cm−1 at room temperature with a relative density of 92%. However, LLZO was found in two crystalline polymorphs: a low-temperature stable tetragonal structure with a space group of I41/acd and a high-temperature stable cubic structure with a space group of Ia3d. Figure 11 shows the tetragonal phase of the garnet-type solid conductor LLZO with its connectivity pattern in the lattice structure. The garnet-type tetragonal LLZO framework structure consists of edge-sharing dodecahedral La(1)O8, La(2)O8 and octahedral ZrO6 units. In the tetragonal framework structure, Li atoms occupy three types of crystallographic sites: Li(1) atoms occupied the tetrahedral-8a sites while the other tetrahedral-16e was unoccupied, Li(2) and Li(3) atoms fully occupied octahedral-16f and octahedral-32g sites, respectively. Awaka et al.67 synthesized a tetragonal Li7La3Zr2O12 garnet by using the flux method; this garnet showed a bulk conductivity of 1.63 × 10−6 S cm−1 and a grainboundary conductivity of 5.59 × 10−7 S cm−1 at room temperature. Wolfenstine et al.68


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suggested that the use of hot-press technique could further enhance the conductivity of tetragonal Li7La3Zr2O12. The hot-pressed tetragonal LLZO showed a high relative density of ~98% and exhibited a total ionic conductivity of ~2.3 × 10−5 S cm−1 at room temperature. Choi et al.69 synthesized composite membranes composed of organic polymer matrixes and inorganic solid electrolytes. The conductivity of as-synthesized tetragonal LLZO pellet-type membrane showed 6.2 × 10−7 S cm−1 at room temperature. The ionic conductivity was synergistically enhanced by the incorporation of the inorganic SE (LLZO) into the organic polymer matrix (PEO-LiClO4). The composite membrane containing 52.5% LLZO displayed an ionic conductivity of 4.42 × 10−4 S cm−1 at 55 °C.69 The Li-ion conductivity of the cubic phase is higher than that of the tetragonal phase by two orders of magnitude for the garnet-type electrolytes. Therefore, to stabilize the cubic phase of garnet, researchers have developed many methods and the

Figure 12. Parameters in obtaining the cubic phase of the garnet structure.


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conductivity was improved to the order of ~10−4 to 10−3 S cm−1 at room temperature (Figure 12). Awaka et al.70 studied the detailed crystal structure of cubic Li7La3Zr2O12 by single-crystal X-ray structure analysis. In the cubic phase, the Li(1) atoms occupied the tetrahedral-24d sites and the Li(2) atoms occupied the distorted octahedral-96h sites. The disordered cubic structure and partial occupation of Li atoms at the Li(2) sites play key roles for fast ion conduction. In the cubic garnet structure, Li-ions could conduct from 24d sites to 96h sites through a three-dimensional pathway. The basic unit of the arrangement is drawn as a loop built by the Li(1) and Li(2) sites (Figure 13a-b). Geiger et al.71 found that LLZO synthesized in Pt crucible led to a tetragonal phase, whereas that synthesized in alumina crucible became a mixture of tetragonal and cubic phases. The accidental incorporation of Al from crucible to the garnet should contribute to stabilize

Figure 13. (a) The loop structure constructed by arrangement of Li atoms, the occupancy value g for each site is noted in the parenthesis. (b) Three-dimensional network structure of the arrangement of Li atoms in cubic Li7La3Zr2O12.70 Reproduced with permission from The Chemical Society of Japan.


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the cubic phase. Rangasamy et al.72 examined the effect of doping of Al into the Li7La3Zr2O12 garnet structure by intentionally incorporating Al during synthesis, and the cubic phase was stabilized by optimizing Al amount to 0.24 in molecular formula. The hot-pressed cubic garnet Li6.24La3Zr2Al0.24O11.98 displayed a conductivity of 4.0 × 10−4 S cm−1 at room temperature with a relative density of 98%. Xu et al.73 proposed a multistep sintering process for the synthesis of the LLZO Li7La3Zr2O12 doped with 0.2 mol% Al2O3 by holding the sample at 900 °C for 6 h and at 1100 °C for 6 h and finally at 1200 °C for 12 h. By holding the sample at 900 °C for 6 h, the precursors undergo partial decomposition to yield the mixture of tetragonal and cubic phase; by continuously holding at 1100 °C for 6 h, the formed tetragonal LLZO completely turns into cubic garnet; finally by holding at 1200 °C for 12 h, the relative density of the sample increases without formation of impurities in cubic phase. They reported that multi-step sintering enhanced the relative density of garnet and the ionic conductivity. The Al-LLZO pellet sintered via a multi-step process exhibited a pure cubic phase with a relative density of 94.25% and displayed an ionic conductivity of 4.5 × 10−4 S cm−1 at room temperature.73 Moreover, 0.25 mol% Al doped Li7La3Zr2O12 samples were synthesized in Pt and alumina crucibles, and ionic conductivity and air stability of the samples were studied.74 The Al-doped LLZO garnet sintered in Pt crucible exhibited higher relative density, conductivity, and air stability than that sintered in alumina crucible. The sample sintered in alumina crucible found with excess Al content coming from alumina crucible may undergo Li loss during high temperature sintering. The LLZO garnet sintered in Pt crucible displayed an ionic conductivity of 4.48 × 10−4 S cm−1 at room temperature with a relative density of ~96%.74 El-Shinawi et al.75 have developed hybrid sol-gel solid-state


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Figure 14. A schematic representation of stabilization and densification of fast-ion conducting LLZO by (a) Interaction with the alumina crucible at sintering temperature 1230 °C (b) with nanostructured alumina at 1100 °C.

approach to synthesize Al-doped Li7La3Zr2O12 with low Al-incorporation level (~0.12 mol%). To obtain dense cubic Li7La3Zr2O12-type phases, however, often required the prolonged calcination at elevated temperatures of up to 1230 °C.66 However, using hybrid sol-gel solid-state approach dense LLZO ceramic pellets obtained by mixing Al2O3 nanosheets with sol-gel processed LLZO solid precursor and followed by single calcination step at lower temperature, 1100 °C for 3 h (Figure 14). The minor secondary phases of LiAlO2 and Li2ZrO3 formed in situ during the calcination process could help to stabilize the cubic phase and increase the density of the garnet. The Al-doped Li7La3Zr2O12 exhibited total conductivity of ~3 × 10−4 S cm−1 at room temperature with an activation energy of 0.27 eV.75 Dumon et al.76 synthesized strontium-doped cubic Li7La3Zr2O12 by adding 0.9‒8.4 wt% Sr. The additive SrCO3 acts as a sintering aid to stabilize the cubic phase, and it increased the grain size and enhanced the conductivity.


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The total ionic conductivity of 1.7 wt% Sr-doped LLZO was approximately 5 × 10−4 S cm−1 at room temperature with an activation energy of approximately 0.31 eV.76 Very recently, our group77 demonstrated that the conductivity of Al-doped LLZO solid electrolyte can be enhanced by employing cheap and relatively simple voltammetric treatment in an all-solid-state Li | LLZO | Li cell configuration. The Li depositiondissolution signal has been observed in the voltammograms, which was supported by neutron powder diffraction measurements, which showed that lithium content in the lattice increased after voltammetric treatment. Furthermore, a local rearrangement of O atoms was detected by X-ray photoelectron spectroscopy, which indicated reduction of oxygen defects. The enhancement in conductivity of LLZO was attributed to both the reduction of oxygen defect and the increase of lithium content. The garnet type Li7La3Zr2O12 shows remarkably high ionic conductivities upon doping of the supervalent Ta, Nb, and Y ions at the Zr sites.78 Ohta et al.11 synthesized the Nb-doped LLZO Li7-xLa3Zr2-xNbxO12 (x = 0 to 2) by partially doping Nb at Zr site; this Li6.75La3Zr1.75Nb0.25O12 displayed a high ionic conductivity of 8 × 10−4 S cm−1 at room temperature. The increase in conductivity upon Nb substitution is due to the structural modification around the lithium sites and increase in Li+ mobility. By doping alkaline earth metal element Ca at the Zr sites, Song et al.79 enhanced the conductivity of LLZO. The doped Ca-ions form octahedral geometry at Zr sites and expand the lattice thus facilitate the migration of Li-ions through bottlenecks with increased size. The Cadoped Li7.1La3Zr1.95Ca0.05O12 exhibited the conductivity of 5.2 × 10−4 S cm−1 at room temperature with an activation energy of 0.27 eV and was electrochemically stable up to 9 V vs. Li/Li+.79 Allen et al.78 investigated the cubic garnet Li6.75La3Zr1.75Ta0.25O12


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(LLZTO) by doping Ta into the LLZO framework; LLZTO displayed a relatively high total ionic conductivity of 8.7 × 10−4 S cm−1 at room temperature with a low activation energy of 0.22 eV. Li et al.80 reported that the ionic conductivity of cubic Li6.75La3Zr1.75Ta0.25O12 (LLZTO-Al) SE prepared in an oxygen sintering atmosphere was approximately 7.4 × 10−4 S cm−1 at room temperature. Buannic et al.81 have enhanced the conductivity of LLZO by doping Ga and Sc at the Li and Zr sites, respectively. The first dopant, Ga3+ doped at Li sites, stabilizes the cubic garnet phase and simultaneously incorporation of second dopant Sc3+ at Zr4+ site, increases amount of mobile Li-ions. The dual substituted LLZO, Li6.65Ga0.15La3Zr1.90Sc0.10O12 exhibited the highest conductivity of 1.8 × 10−3 S cm−1 at room temperature with an activation energy of 0.29 eV. Gu et al.82 studied the influence of pentavalent and trivalent doping on LLZO garnet by combining experimental results with molecular dynamic simulations. They showed that the conductivity of cubic garnet was enhanced upon pentavalent doping (Ta5+ or Nb5+) which increased the disorder and vacancy concentrations, whereas the trivalent doping of Al3+ or Ga3+ on the Li sites was showed to be slightly less effective. The immobile Al3+ or Ga3+ preferentially occupies 24d sites which blocks conduction pathway of Li-ions. The doping of Ga3+ on the La sites stabilized the cubic phase and had no effect on the Li+ concentration.82 Parameters in obtaining the cubic phase of the garnet structure are shown in Figure 12, including the sintering temperature and time, additives, sintering atmosphere, element doping, Li content (occupation sites), crucibles etc. Bernstein et al.83 studied the driving force behind the phase transition from low conducting tetragonal phase to high conducting cubic phase by using density-functional theory and a variable cell shape version of molecular dynamics. DFT calculations showed


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that the Li sublattice in tetragonal phase was ordered which was either fully occupied or unoccupied with Li-ions, whereas that in cubic phase was disordered which was partially filled with Li-ions. In the cubic structure, the tetrahedral 24d sites (Li1c) were fully occupied and surrounded by four pairs of octahedral 96h sites (Li2c), which were partially occupied. The Li2-Li2 pair was energetically prohibited to occupy both sites at the same time due to the short distance between the pair, which causes strong Coulomb repulsion between two Li-ions. In the tetragonal structure, cubic Li1c sites converted to fully occupied 8a sites (Li1t) and unoccupied 16e sites (Li2ut). The partially filled cubic Li2c sites adapted into fully occupied 16f (Li2t) and 32g (Li3t) sites in tetragonal structure (Figure 15). For MD simulations, 0.25 vacancies per formula unit were applied at 600 K.

Figure 15. Li sublattice in the cubic (left) and tetragonal (right) phases of LLZO. All Li positions are included, although in the cubic phase not all are occupied. The Li(1) atoms (8at and 24dc) are large gray (gold), Li(2) atoms (16ft and 96hc) are white, and Li(3) atoms (32gt) are dark gray. The cubic Li(1) positions that become vacant upon transition to the ordered tetragonal structure (16et) are indicated by small (gold) spheres.83 Reproduced with permission from American Physical Society.


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The vacancies were formed randomly in the ordered, tetragonal unit cell by removing Liions from Li1t and Li3t sites. The system spontaneously transformed into cubic within 510 ps of simulation time. At 15 ps, the system fluctuated between cubic and tetragonal states, and finally settled into a cubic phase after 30 ps. When the tetragonal phase transforms into cubic phase, the ratios of lattice constants, ax/az and ay/az (the lattice constants along x and y to that along z) drop from 1.04 to 0.98 (Figure16, top panel). When this drop occurs, there is a re-distribution of Li-ions in tetrahedral sites. The occupancies of tetrahedral Li1t (red) sites and tetrahedral Li2ut (blue) sites are initially near 1 and 0, respectively, in consistent with experimental tetragonal structure (Figure16,

Figure 16. Evolution over time of structure and site occupation quantities for a sample system with nvac = 0.25 (vacancy number) at T = 600 K. Top: unit cell shape (ax/az blue, ay/az red) and volume (black). Middle: 96hc (black) and 16ft+32gt (red) lattice site occupations. Bottom: 24dc (black), 8at (red) and 16et (blue with symbols) lattice site occupations.83 Reproduced with permission from American Physical Society.


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bottom panel). When the system transforms from tetragonal to cubic, there is a sharp drop in the occupancies of Li1t (red, bottom panel), Li2t and Li3t sites (red, middle panel), while Li2ut sites (blue, bottom panel) observed sharp increment in their occupancies. However, the cubic Li1c (black, bottom panel) and Li2c (black, middle panel) sites did not experience increase or decrease in their occupancies during phase transition, indicating that Li-ion re-distribution took place between occupied Li1t sites and unoccupied Li2ut sites during simulation studies. So the disorder in cubic structure calculated originated not only from the change in the lattice constants, but also from

Figure 17. The blocking effect of Al in cubic LLZO. Simulated BVM data is shown as blue shaded area. (a) 3D illustration of un-doped cubic LLZO. Projected view from along the [100] direction of (b) un-doped cubic LLZO, (c) Al-doped LLZO with Al in 24 d sites, (d) Al-doped LLZO with Al in 96 h sites. The paths blocked by Al are marked with red “X” in (c) and (d). The green arrows are depicted to represent diffusive motion of Li.84


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Figure 18. The effect of additional Ta doping on Al site preference in Al-doped LLZO. The calculated energies of LLZO with Al in 24d and 96h site are described as ∆ and □, respectively. All calculated energies are rescaled by making the most stable configuration in each doping group correspond to zero. Inset image describes two different doping sites of Al.84

filling up the empty Li2ut octahedral sites. Shin et al.84 have studied the enhancement in conductivity and lithium-ion transport phenomena in cubic Li7La3Zr2O12 upon multi-doping by combining the experimental and density functional theory calculations. The blocking effect of immobile Al3+ at 24d sites and 96h sites in cubic structure is shown in Figure 17. When Al occupies 24d sites, the relative energy significantly decreased by 11.9 meV per atom, compared with the condition that Al occupies 96h sites, suggesting that Al occupation at 24d sites is more favored (Figure 18). It is widely believed that the immobile Al dopant in the Li sites may reduce the Li-ion diffusion by blocking the conduction paths of Li ions in the


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garnet.78 The blocking effect of Al at 24d sites is more serious in decreasing the ionic conductivity compared with the blocking effect of Al at 96h sites. By multi-doping on LLZO garnet with Al and Ta atoms, the 96h sites were stabilized and the preference of Al occupation at 24d sites were reduced. At higher level of Ta doping, the 96h sites were further stabilized, indicating that the blocking of Al at 24d junction sites was reduced by multi-doping

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LLZO.

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multi-doped

cubic

Li6.2Al0.2La3Zr1.8Ta0.2O12 exhibited about 6.14 × 10−4 S cm−1 at room temperature with activation energy of 0.29 eV. Kotobuki et al.85 successfully used the lithium-ion conducting garnet Li7La3Zr2O12 (LLZO) as electrolytes to fabricate all-solid-state rechargeable lithium batteries with Li metal. The cyclic voltammogram of the Li/LLZO/Li cell shows that the dissolution and deposition reactions of lithium occurred reversibly without any reaction with LLZO, indicating stability of the LLZO against Li metal. A full cell with a LiCoO2/LLZO/Li configuration was successfully operated and demonstrated a discharge capacity of 15 µA h cm−2. However, an irreversible behavior was observed at the first charge and discharge cycle due to an interfacial issue between LiCoO2 and LLZO.85 Kim et al.86 reported that the interfacial layer La2CoO4 formed between the Li+ conducting SE LLZO and LiCoO2 cathode due to mutual diffusion, suppressing the lithium insertion/extraction at LLZO/LiCoO2 interface.86 Ohta et al.87 fabricated an all-solid-state lithium-ion battery consisting of LiCoO2/Li6.75La3Zr1.75Nb0.25O12/Li configuration to investigate its electrochemical performance and charge transfer resistance. The cell showed good cycle performance and displayed a capacity retention of approximately 98% after 100 charge-discharge cycles. The discharge capacity at the first cycle was 129


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mAh g−1 and stabilized at 127 mAh g−1 after the 100th cycle (theoretical capacity of LCO is 137 mAh g−1, which corresponds to 0.5 Li per CoO2).87 Moreover, they introduced Li3BO3 as buffer layer between LiCoO2 cathode and Nb-doped LLZO SE to increase the interfacial contact between them. The obtained battery showed a good charge-discharge capacity with a low interfacial resistance between electrodes and SE. The charge and discharge capacities at the first cycle were 100 and 85 mAh g−1, respectively (theoretical capacity of LCO is 115 mAh g−1, which corresponds to 0.42 Li per CoO2).88 Du et al.89 fabricated all-solid-state Li batteries using Ta-doped LLZO coated with a composite cathode which is a mixture of PVDF:LiTFSI, Ketjen Black and carbon-coated LiFePO4 on one side of the pellet, and attached to a Li anode on the other side. The first discharge capacity of the battery was 150 mA h g−1 at 0.05 C and showed a 93% capacity retention after 100 cycles at 60 °C. The performance of the battery can be further improved by increasing the temperature up to 100 °C.89 Jin et al.90 have studied the stability of Aldoped Li7La3Zr2O12 in air by fabricating a Li/Li7La3Zr2O12/Cu0.1V2O5 solid-state battery. The Al-doped Li7La3Zr2O12 garnet was found to be unstable in humid air due to the formation of impurity phases like La(OH)3 and LiOH·H2O, indicating the decomposition of Al-doped Li7La3Zr2O12. The increase in pH of LLZO powder/water suspension was resulted from H+/Li+ exchange reaction between water and the LLZO electrolyte. The conductivity of as-synthesized LLZO was declined from 2.4 × 10−4 S cm−1 to 1.6 × 10−4 S cm−1 at room temperature after one week exposure to air. Ahn et al.91 also examined the effect of moisture on the Li-ion conduction of Li7La3Zr2O12 garnet by composing an allsolid-state battery with a LiFePO4 (LFP) film as cathode. The severe degradation of Liion conduction was observed when the specimen was exposed to humid air due to the


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formation of secondary phase (La(OH)3) at the grain boundary. The charge and discharge capacities of the battery were observed to be extremely low (6 × 10−4 µAh cm−2) at room temperature due to the reduced Li-ion conduction and the capacity was improved with increase of temperature. Yan et al.92 fabricated an all-solid-state battery composed of Li/Li7La3Zr2O12/LiFePO4 using an ultrathin LLZO solid electrolyte film. The solid electrolyte film with thickness of several micrometers was prepared by mixing LLZO nanoparticle slurry with appropriate solvent, dispersant, adhesives and surfactant without cold or hot-pressing. The Li/LLZO/LFPO cell exhibited the first discharge capacity of 160.4 mAh g−1 at room temperature and retained capacity of 136.8 mAh g−1 after 100 cycles. Thangadurai et al.93 published a critical review on the garnet-type solid-state fast Li-ion conductors for Li batteries.

Perovskite-type Li-ion conductors: The perovskite-type Li3xLa(2/3)-xTiO3 (LLTO, 0 < x Li-excess LLTO (0.391 eV) (Figure 19e). Chen et al.101 have reported the improved ionic conductivity of Li0.35La0.55TiO3 by mixing Li7La3Zr2O12 (LLZO) sol into its precursor powder. The Zr doped into the LLTO mainly prefers the


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Figure 19. (a) A schematic of crystal structures of LLTOs; a tetragonal structure (left) and an orthorhombic structure (right). FE-SEM micrographs of LLTOs sintered at (b) Low-T and (c) High-T (d) A comparison of Li+ conductivities along with schematics of the domain microstructures. (e) Arrhenius plots of the boundary conductivities for low-T LLTO, high-T LLTO and Li-excess LLTO.100 Reproduced with permission from Royal Society of Chemistry.

grain boundary region. The grain boundary conductivity of LLTO was largely increased


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by the introduction of LLZO. The LLTO-based electrolyte mixed with 5 wt.% LLZO exhibited a total conductivity of 1.2 × 10−4 S cm−1 and a grain boundary conductivity of 1.5 × 10−4 S cm−1 at room temperature. The perovskite-type Li0.35La0.55TiO3 (LLTO) with high ionic conductivities of 10−3 - 10−4 S cm−1 at room temperature could be used as the electrolytes for all-solid-state batteries (ASSBs). However, the reduction of Ti4+ by Li metal limits their applicability in ASSBs. Kotobuki et al.102 have fabricated an all-solid-state battery with the composition of LiCoO2/LLT/Li4Mn5O12 using a honeycomb-structured Li0.35La0.55TiO3, which has micro-sized holes on both sides of the membrane. The impregnation of active materials LiCoO2 and Li4Mn5O12 particles mixed with respective precursor sols into the honeycomb holes forms good active materials/solid electrolyte interfaces, which reduces the internal resistance of the cell and thus improves the discharge capacity. Li et al.103 have

showed

the

charge/discharge

performance

of

an

all-solid-state

Li/LiPON/LLTO/LiCoO2 cell using LiPON/LLTO thin-film as solid electrolyte. Amorphous lithium lanthanum titanate (LLTO) solid electrolyte thin-films have been fabricated using e-beam evaporation. LiPON thin-film was used as coating layer to avoid undesirable reaction between LLTO and lithium metal. The cell exhibited a discharge capacity of about 50 µA h/cm2 and the capacity degradation was about 0.5% per cycle after 100 charge-discharge cycles at a current of 7 µA/cm2. Le et al.104 have examined the stability and durability of Al-substituted LLTO ceramic by immersing it in aqueoussolution for one month. The Al-substituted lithium lanthanum titanate (A-LLTO) sintered at 1350 °C for 6 h exhibited ionic conductivity of about 3.17 × 10−4 S cm−1 at room temperature. The Li-LiCoO2 and Li-O2 cells with LiPON/A-LLTO as electrolytes


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exhibited good cyclability, stability and superior electrochemical performance after the electrolytes being immersed in an alkaline aqueous solution. The charge-discharge capacity of the Li-LiCoO2 cell maintained 59.3% of the initial capacity with a coulombic efficiency of 98.3 % after 100 cycles at 1C rate. Two comprehensive reviews on the structure and properties of LLTO were published by Bohnke96 and Stramare et al.105.

LISICON-type Li-ion conductors: A polycrystalline Li+ conducting SE LISICON Li2+2xZn1-xGeO4 was first reported by Bruce and West in 1983.106 The LISICON framework is similar to the γ-Li3PO4 crystal structure. Although Li3.5Zn0.25GeO4 stoichiometry (x = 0.75) showed the highest lithiumion conductivity (0.125 S cm−1) at high temperature (300 °C), its conductivity at room temperature is only 1 × 10−7 S cm−1. For LISICON compounds, the ionic conductivity tends to decrease with time at low temperature due to the formation of a defective complex, namely, Li4GeO4, which traps the mobile lithium-ions through the immobile sub-lattice.107 Attempt to restore their conductivities to original values by re-annealing of the samples was unsuccessful in Li2+2xZn1-xGeO4 compounds due to aging problem. Furthermore, LISICON, Li14ZnGe4O16 is highly reactive with lithium and atmospheric CO2.108 In the LISICON family, various materials displaying the γ-Li3PO4 framework have been synthesized. The germanium-doped framework (Li3.6Ge0.6V0.4O4) displayed the highest conductivity of approximate 4 × 10−5 S cm−1 at room temperature.109 Deng et al.110 reported enhanced conductivities of LISICON SEs by doping Si in Li4SiO4 system with P, Al or Ge. They have synthesized a series of compounds, including Li3.75Si0.75P0.25O4, Li4.25Si0.75Al0.25O4, Li4Al0.33Si0.33P0.33O4 and Li4Al0.33Si0.17Ge0.17P0.33O4


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Figure 20. Three types of Li-ion diffusion mechanisms for substituted Li4SiO4 NASICONs. Li-ions are shown in yellow; XO4 groups are shown in red.110 Copyright 2017 American Chemical Society.

with cationic substitution of Si in the parent LISICON-like Li4SiO4 framework. The ionic conductivity of Li4SiO4 is very low (σ473K = 4 × 10−6 S cm−1), upon substitution of P or Al at Si sites the conductivity was increased by 2 orders of magnitude to 1 × 10−4 S cm−1 for Li3.75Si0.75P0.25O4 and 2 × 10−4 S cm−1 for Li4.25Si0.75Al0.25O4 at 473 K. A further increase

in

the

ionic

conductivity

was

observed

for

'ternary'

composition

Li4Al0.33Si0.33P0.33O4 with σ473K = 1 × 10−3 S cm−1. Molecular dynamics simulations (MD) were used to explain the enhanced conductivities and Li-ion diffusion mechanisms for substituted compositions. Depending on the temperature, they have proposed three Li-ion diffusion mechanisms: (i) local oscillation at low temperature, (ii) isolated hopping at intermediate temperature and (iii) superionic motion at high temperature (Figure 20). MD simulations reveal that the Li-ion diffusion mechanism in each composition was temperature-dependent. The Li-ion mobility became higher when the diffusion mechanism transform from local oscillation to isolated hopping or from isolated hopping to superionic flow. The mixed polyanion substitution lowers the temperature at which the


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transition to the superionic state with high Li-ion conductivity occurs.110 Song et al.111 have reported a facile strategy by substituting Cl for O to adjust bottleneck size and binding energy to enhance the conductivities and electrochemical stability of LISICON systems. The lattice parameters of Cl-substituted Li10.5-xSi1.5P1.5ClxO12-x solid solutions are larger than that of the pure Li3PO4, indicating expansion of bottlenecks and decrease of energy barriers, which enhanced the ionic conductivities. Li10.42Si1.5P1.5Cl0.08O11.92 and Li10.42Ge1.5P1.5Cl0.08O11.92 exhibited ionic conductivities of about 1.03 × 10−5 S cm−1 and 3.7 × 10−5 S cm−1 at room temperature, respectively, with electrochemical stability windows of up to 9 V vs. Li/Li+. All-solid-state batteries were assembled with 0.3Li2MnO3·0.7LiMn1.5Ni0.5O4 as cathode and lithium metal as anode. The battery based on the Li10.42Si1.5P1.5Cl0.08O11.92 showed initial charge and discharge capacities of 114.5 mA h g−1 and 114.1 mA h g−1, respectively, with coulombic efficiencies of 100%. The battery based on the Li10.42Ge1.5P1.5Cl0.08O11.92 showed initial charge and discharge capacities of 330.1 mA h g−1 and 133.2 mA h g−1, respectively. The battery based on Li10.42Si1.5P1.5Cl0.08O11.92 showed better cycle performance than that based on Li10.42Ge1.5P1.5Cl0.08O11.92.111

Interfacial stability: The use of highly conductive solid-state inorganic electrolytes in all-solid-state lithiumion batteries has been extensively researched in recent years because of the safety and good performance of this state-of-the-art battery technology. However, incorporation of these materials in all-solid-state batteries remains very challenging due to their reactivity with the electrode materials at interfaces. Introduction of a buffer layer between the


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electrode and electrolyte reduces interfacial resistance. Kato et al.112 introduced a thin Nb layer (~10 nm) to interface between SE Li7La3Zr2O12 and cathode LiCoO2 to reduce interfacial resistance. The Nb layer suppresses the growth of mutual diffusion layer at the interface and produces an amorphous Li-Nb-O material through an in-situ process, which is reported to be Li+ conductive. Introduction of the Nb layer reduces the interfacial resistance and improves the cycle stability and rate capability of charge-discharge reactions of a Li/Li7La3Zr2O12/LiCoO2 solid-state battery.112 Zhou et al.113 constructed an all-solid-state battery using a NASICON Li1.3Al0.3Ti1.7(PO4)3 (LATP) ceramic membrane sandwiched with a crosslinked polymer, poly(ethylene glycol) methyl ether acrylate (CPMEA), on both sides. Insertion of polymer layer between SE and anode suppresses dendrite formation due to uniform Li+ flux across the polymer/lithium interface. At the same time, the polymer layer showed a better ability of wetting toward lithium metal, and the ceramic layer was protected from contact with lithium metal. The all-solid-state battery, Li/CPMEA/LATP/CPMEA/LiFePO4 cell, delivered a stable capacity of 130 mAh g−1 after 100 cycles with coulombic efficiency of 99.7%–100%.113 Garnet-type LLZO SEs could react with carbon dioxide and moisture in ambient air to form Li-ioninsulating Li2CO3 layers on the surfaces, resulting in large interfacial resistances. Moreover, they reported that by addition of 2 wt% LIF to the garnet Li6.5La3Zr1.5Ta0.5O12, stability of the garnet electrolyte against moist air was effectively increased. Additionally, they fabricated an all-solid-state battery, Li/polymer/LLZT-2LiF/LiFePO4, which demonstrated high coulombic efficiency and long cycle life with reduced interfacial resistance. Due to the low interfacial resistance, the low overpotential further decreased to 0.2 V at 80 µA cm−2.114


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Summary and Future Outlook In conclusion, LIBs with liquid electrolytes are still providing a large part of the current energy demand. However, serious safety concerns arise due to the usage of liquid electrolytes. Leakage of electrolytes and growth of Li-dendrites caused by contact of anodes with liquid electrolytes greatly hinder the performance of LIBs. Devices operated at elevated temperatures also suffer from electrolyte evaporation while liquid electrolytes are used. Hence, inorganic SEs with room temperature conductivities of >10−3 S cm−1 and wide electrochemical stability windows (6 V vs. Li/Li+) have been intensively studied to develop all-solid-state batteries that can be used for applications with wide temperature range. Stability of SE/electrode interface is a challenge for facile passage of Li-ions across the interface. From this perspective, we have discussed recent published research activities in synthetic methods and possible dopants, and respective effects on the crystal lattices of NASICON-type, garnet-type, perovskite-type and LISICON-type SEs. Doping with isovalent or aliovalent elements and addition of additives into the parent structural frameworks enhance the ionic conductivities of SEs due to the framework expansion and increase of mobile Li-ions. We have also discussed the interfacial reactions between SEs and electrodes. For example, the solid Li-ion conductor LiTi2(PO4)3-based material, which exhibited high ionic conductivity, was chemically unstable with Li metal due to the reduction of Ti4+. Introduction of a protective layer (LiPON film) between SE and Li metal improved the stability. Good interfacial contacts between the solid electrolytes and the electrodes can be achieved given the nondeformable nature of the inorganic ceramic SEs. However, further studies are needed to better understand the interfacial processes between solid electrolytes and electrodes.


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ASSOCIATED CONTENT Biography: Yedukondalu Meesala has obtained PhD in chemistry from Indian Institute of Technology, Bombay, India (2011). He is currently a post-doctoral fellow in Prof. Ru-Shi Liu’s group at the Department of Chemistry, National Taiwan University, Taiwan. His current research interest focuses on transport properties of solid electrolyte materials for Liion batteries. Anirudha Jena has received PhD degree from Indian Institute of Science, Bengaluru, India (2013). He is currently working as a Research Assistant Professor in the Graduate Institute of Manufacturing Technology, National Taipei University of Technology. His current research is focused on the development of materials for all-solid-state batteries. Ho Chang has received PhD degree from the National Taipei University of Technology (NTUT) (2004). He is currently working as a professor at the Graduate Institute of Manufacturing Technology, NTUT. His current research is focused on dynamic frication behaviors of pneumatic cylinders, fabrication of thermoelectric materials. Ru-Shi Liu obtained two PhD degrees in chemistry - one from the National Tsing Hua University in 1990 and another from the University of Cambridge in 1992. He is currently working as a professor at the Department of Chemistry, National Taiwan University. His research interest is materials chemistry. He is an author and coauthor of more than 530 publications in international scientific journals.

Notes The authors declare no competing financial interests. ACKNOWLEDGMENT This study was financially supported by the Ministry of Science and Technology, Taiwan (Contract No: MOST 104-2113-M-002-012-MY3). REFERENCES 1. Vetter, J.; Novák, P.; Wagner, M. R.; Veit, C.; Möller, K. C.; Besenhard, J. O.; Winter, M.; Wohlfahrt-Mehrens, M.; Vogler, C.; Hammouche, A. Ageing Mechanisms in Lithium-Ion Batteries. J. Power Sources 2005, 147, 269-281. 2. Tarascon, J.-M.; Armand, M. Issues and Challenges Facing Rechargeable Lithium Batteries. Nature 2001, 414, 359-367.


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3. Gachot, G.; Grugeon, S.; Armand, M.; Pilard, S.; Guenot, P.; Tarascon, J.-M.; Laruelle, S. Deciphering the Multi-Step Degradation Mechanisms of Carbonate-Based Electrolyte in Li Batteries. J. Power Sources 2008, 178, 409-421. 4. Xu, W.; Wang, J.; Ding, F.; Chen, X.; Nasybulin, E.; Zhang, Y.; Zhang, J.-G. Lithium Metal Anodes for Rechargeable Batteries. Energy Environ. Sci. 2014, 7, 513537. 5. Cheng, X. B.; Zhang, R.; Zhao, C. Z.; Wei, F.; Zhang, J. G.; Zhang, Q. A Review of Solid Electrolyte Interphases on Lithium Metal Anode. Adv. Sci. 2016, 3, 1500213. 6. Aricò, A. S.; Bruce, P.; Scrosati, B.; Tarascon, J.-M.; Van Schalkwijk, W. Nanostructured Materials for Advanced Energy Conversion and Storage Devices. Nat. Mater. 2005, 4, 366-377. 7.

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18. Xu, H.; Wang, S.; Wilson, H.; Zhao, F.; Manthiram, A. Y-Doped NASICON-type LiZr2(PO4)3 Solid Electrolytes for Lithium-Metal Batteries. Chem. Mater. 2017, 29, 7206-7212. 19. Xiong, L.; Ren, Z.; Xu, Y.; Mao, S.; Lei, P.; Sun, M. LiF Assisted Synthesis of LiTi2(PO4)3 Solid Electrolyte with Enhanced Ionic Conductivity. Solid State Ionics 2017, 309, 22-26. 20. Sljukic, M.; Matkovic, B.; Prodic, B.; Scavnicar, S. Preparation and Crystallographic Data of Phosphates with Common Formula MiM42(PO4)3 (Mi = Li Na K Rb Cs-M4 = Zr Hf). Croat. Chem. Acta. 1967, 39, 145. 21. Sudreau, F.; Petit, D.; Boilot, J. Dimorphism, Phase Transitions, and Transport Properties in LiZr2(PO4)3. J. Solid State Chem. 1989, 83, 78-90. 22. Subramanian, M.; Subramanian, R.; Clearfield, A. Lithium Ion Conductors in the System AB(IV)2(PO4)3 (B = Ti, Zr and Hf). Solid State Ionics 1986, 18, 562-569. 23. Aono, H.; Sugimoto, E.; Sadaaka, Y.; Imanaka, N.; Adachi, G. Ionic Conductivity of the Lithium Titanium Phosphate (Li1+xAlxTi2-x(PO4)3, M = Al, Sc, Y, and La) Systems. J. Electrochem. Soc. 1989, 136, 590-591. 24. Francisco, B. E.; Stoldt, C. R.; M’Peko, J.-C. Lithium-Ion Trapping from Local Structural Distortions in Sodium Super Ionic Conductor (NASICON) Electrolytes. Chem. Mater. 2014, 26, 4741-4749. 25. Kwatek, K.; Nowiński, J. L. Electrical Properties of LiTi2(PO4)3 and Li1.3Al0.3Ti1.7(PO4)3 Solid Electrolytes Containing Ionic Liquid. Solid State Ionics 2017, 302, 54-60. 26. Li, S.-c.; Cai, J.-y.; Lin, Z.-x. Phase Relationships and Electrical Conductivity of Li1+xGe2-xAlxP3O12 and Li1+xGe2-xCrxP3O12 Systems. Solid State Ionics 1988, 28, 12651270. 27. Aono, H.; Sugimoto, E.; Sadaoka, Y.; Imanaka, N.; Adachi, G.-y. Electrical Property and Sinterability of LiTi2(PO4)3 Mixed with Lithium Salt (Li3PO4 or Li3BO3). Solid State Ionics 1991, 47, 257-264. 28. Arbi, K.; Bucheli, W.; Jiménez, R.; Sanz, J. High Lithium Ion Conducting Solid Electrolytes Based on NASICON Li1+xAlxM2-x(PO4)3 Materials (M = Ti, Ge and 0 ≤ x ≤ 0.5). J. Eur. Ceram. Soc. 2015, 35, 1477-1484. 29. Bucheli, W.; Jimenez, R.; Sanz, J.; Várez, A. The Log (σ) Vs. Log (ω) Derivative Plot Used to Analyze the Ac Conductivity. Application to Fast Li+ Ion Conductors with Perovskite Structure. Solid State Ionics 2012, 227, 113-118. 30. Shang, X.; Nemori, H.; Mitsuoka, S.; Mastuda, Y.; Takeda, Y.; Yamamoto, O.; Imanishi, N. High Lithium-Ion Conducting NASICON-Type Li1+x-yAlxNbyTi2-x-y(PO4)3 Solid Electrolytes. Solid State Ionics 2016, 297, 43-48. 31. Shimonishi, Y.; Zhang, T.; Imanishi, N.; Im, D.; Lee, D. J.; Hirano, A.; Takeda, Y.; Yamamoto, O.; Sammes, N. A Study on Lithium/Air Secondary Batteries—Stability of the NASICON-Type Lithium Ion Conducting Solid Electrolyte in Alkaline Aqueous Solutions. J. Power Sources 2011, 196, 5128-5132.


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32. Kothari, D. H.; Kanchan, D. K. Study of Study of Electrical Properties of Gallium-Doped Lithium Titanium Aluminum Phosphate Compounds. Ionics 2014, 21, 1253-1259. 33. Leo, C.; Rao, G. S.; Chowdari, B. Effect of MgO addition on the ionic conductivity of LiGe2(PO4)3 ceramics. Solid State Ionics 2003, 159, 357-367. 34. Chung, H.; Kang, B. Increase in Grain Boundary Ionic Conductivity of Li1.5Al0.5Ge1.5(PO4)3 by Adding Excess Lithium. Solid State Ionics 2014, 263, 125-130. 35. Fu, J. Superionic Conductivity of Glass-Ceramics in the System Li2O-Al2O3TiO2-P2O5. Solid State Ionics 1997, 96, 195-200. 36. Fu, J. Fast Li+ Ion Conducting Glass-Ceramics in the System Li2O-Al2O3-GeO2P2O5. Solid State Ionics 1997, 104, 191-194. 37. Chowdari, B.; Rao, G. S.; Lee, G. XPS and Ionic Conductivity Studies on Li2OAl2O3-(TiO2 or GeO2)-P2O5 Glass-Ceramics. Solid State Ionics 2000, 136, 1067-1075. 38. Thokchom, J. S.; Kumar, B. Water Durable Lithium Ion Conducting Composite Membranes from the Li2O-Al2O3-TiO2-P2O5 Glass-Ceramic. J. Electrochem. Soc. 2007, 154, A331-A336. 39. Mohaghegh, E.; Nemati, A.; Eftekhari Yekta, B.; Banijamali, S. Effects of Fe2O3 Content on Ionic Conductivity of Li2O-TiO2-P2O5 Glasses and Glass-Ceramics. Mater. Chem. Phys. 2017, 190, 8-16. 40. Kumar, B.; Nellutla, S.; Thokchom, J. S.; Chen, C. Ionic Conduction through Heterogeneous Solids: Delineation of the Blocking and Space Charge Effects. J. Power Sources 2006, 160, 1329-1335. 41. Xu, X.; Wen, Z.; Wu, X.; Yang, X.; Gu, Z. Lithium Ion-Conducting GlassCeramics of Li1.5Al0.5Ge1.5(PO4)3-xLi2O (x = 0.0-0.20) with Good Electrical and Electrochemical Properties. J. Am. Ceram. Soc. 2007, 90, 2802-2806. 42. Jadhav, H. S.; Kalubarme, R. S.; Jang, S. Y.; Jung, K. N.; Shin, K. H.; Park, C. J. B2O3-Added Lithium Aluminium Germanium Phosphate Solid Electrolyte for Li-O2 Rechargeable Batteries. Dalton Trans. 2014, 43, 11723-11727. 43. Zhu, Y.; Zhang, Y.; Lu, L. Influence of Crystallization Temperature on Ionic Conductivity of Lithium Aluminum Germanium Phosphate Glass-Ceramic. J. Power Sources 2015, 290, 123-129. 44. Thokchom, J. S.; Kumar, B. The Effects of Crystallization Parameters on the Ionic Conductivity of a Lithium Aluminum Germanium Phosphate Glass-Ceramic. J. Power Sources 2010, 195, 2870-2876. 45. Santagneli, S. H.; Baldacim, H. V. A.; Ribeiro, S. J. L.; Kundu, S.; Rodrigues, A. C. M.; Doerenkamp, C.; Eckert, H. Preparation, Structural Characterization, and Electrical Conductivity of Highly Ion-Conducting Glasses and Glass Ceramics in the System Li1+xAlxSnyGe2-(x+y)(PO4)3. J. Phys. Chem. C 2016, 120, 14556-14567. 46. McMillan, P. W. The Crystallisation of Glasses. J. Non-Cryst. Solids 1982, 52, 67-76.


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47. Xu, X.; Wen, Z.; Yang, X.; Chen, L. Dense Nanostructured Solid Electrolyte with High Li-Ion Conductivity by Spark Plasma Sintering Technique. Mater. Res. Bull. 2008, 43, 2334-2341. 48. Zhang, P.; Matsui, M.; Takeda, Y.; Yamamoto, O.; Imanishi, N. Water-Stable Lithium Ion Conducting Solid Electrolyte of Iron and Aluminum Doped NASICON-Type LiTi2(PO4)3. Solid State Ionics 2014, 263, 27-32. 49. Zhang, P.; Wang, H.; Si, Q.; Matsui, M.; Takeda, Y.; Yamamoto, O.; Imanishi, N. High Lithium Ion Conductivity Solid Electrolyte of Chromium and Aluminum Co-Doped NASICON-Type LiTi2(PO4)3. Solid State Ionics 2015, 272, 101-106. 50. Zhang, M.; Takahashi, K.; Imanishi, N.; Takeda, Y.; Yamamoto, O.; Chi, B.; Pu, J.; Li, J. Preparation and Electrochemical Properties of Li1+ xAlxGe2-x(PO4)3 Synthesized by a Sol-Gel Method. J. Electrochem. Soc. 2012, 159, A1114-A1119. 51. Zhang, M.; Takahashi, K.; Uechi, I.; Takeda, Y.; Yamamoto, O.; Im, D.; Lee, D.J.; Chi, B.; Pu, J.; Li, J. et al. Water-Stable Lithium Anode with Li1.4Al0.4Ge1.6(PO4)3TiO2 Sheet Prepared by Tape Casting Method for Lithium-Air Batteries. J. Power Sources 2013, 235, 117-121. 52. Mariappan, C. R.; Yada, C.; Rosciano, F.; Roling, B. Correlation between MicroStructural Properties and Ionic Conductivity of Li1.5Al0.5Ge1.5(PO4)3 Ceramics. J. Power Sources 2011, 196, 6456-6464. 53. Lang, B.; Ziebarth, B.; Elsässer, C. Lithium Ion Conduction in LiTi2(PO4)3 and Related Compounds Based on the NASICON Structure: A First-Principles Study. Chem. Mater. 2015, 27, 5040-5048. 54. Kang, J.; Chung, H.; Doh, C.; Kang, B.; Han, B. Integrated Study of First Principles Calculations and Experimental Measurements for Li-Ionic Conductivity in AlDoped Solid-State LiGe2(PO4)3 Electrolyte. J. Power Sources 2015, 293, 11-16. 55. West, W. C.; Whitacre, J. F.; Lim, J. R. Chemical Stability Enhancement of Lithium Conducting Solid Electrolyte Plates Using Sputtered LiPON Thin Films. J. Power Sources 2004, 126, 134-138. 56. Jadhav, H. S.; Kalubarme, R. S.; Jadhav, A. H.; Seo, J. G. Highly Stable Bilayer of LiPON and B2O3 added Li1.5Al0.5Ge1.5(PO4)3 Solid Electrolytes for Non-Aqueous Rechargeable Li-O2 Batteries. Electrochim. Acta 2016, 199, 126-132. 57. Jung, Y.-C.; Lee, S.-M.; Choi, J.-H.; Jang, S. S.; Kim, D.-W. All Solid-State Lithium Batteries Assembled with Hybrid Solid Electrolytes. J. Electrochem. Soc. 2015, 162, A704-A710. 58. Safanama, D.; Damiano, D.; Rao, R. P.; Adams, S. Lithium Conducting Solid Electrolyte Li1+xAlxGe2-x(PO4)3 Membrane for Aqueous Lithium Air Battery. Solid State Ionics 2014, 262, 211-215. 59. Safanama, D.; Adams, S. High Efficiency Aqueous and Hybrid Lithium-Air Batteries Enabled by Li1.5Al0.5Ge1.5(PO4)3 Ceramic Anode-Protecting Membranes. J. Power Sources 2017, 340, 294-301.


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60. Liu, Y.; Li, B.; Kitaura, H.; Zhang, X.; Han, M.; He, P.; Zhou, H. Fabrication and Performance of All-Solid-State Li-Air Battery with SWCNTs/LAGP Cathode. ACS Appl. Mater. Interfaces 2015, 7, 17307-17310. 61. Hyooma, H.; Hayashi, K. Crystal structures of La3Li5M2O12 (M = Nb, Ta). Mater. Res. Bull. 1988, 23, 1399-1407. 62. Thangadurai, V.; Weppner, W. Li6ALa2Ta2O12 (A = Sr, Ba): Novel Garnet-Like Oxides for Fast Lithium Ion Conduction. Adv. Funct. Mater. 2005, 15, 107-112. 63. Thangadurai, V.; Kaack, H.; Weppner, W. J. F. Novel Fast Lithium Ion Conduction in Garnet-Type Li5La3M2O12 (M = Nb, Ta). J. Am. Ceram. Soc. 2003, 86, 437-440. 64. Thangadurai, V.; Weppner, W. Effect of Sintering on the Ionic Conductivity of Garnet-Related Structure Li5La3Nb2O12 and In- and K-Doped Li5La3Nb2O12. J. Solid State Chem. 2006, 179, 974-984. 65. Narayanan, S.; Ramezanipour, F.; Thangadurai, V. Enhancing Li Ion Conductivity of Garnet-Type Li5La3Nb2O12 by Y- and Li-Codoping: Synthesis, Structure, Chemical Stability, and Transport Properties. J. Phys. Chem. C 2012, 116, 20154-20162. 66. Murugan, R.; Thangadurai, V.; Weppner, W. Fast Lithium Ion Conduction in Garnet-Type Li7La3Zr2O12. Angew. Chem., Int. Ed. Engl. 2007, 46, 7778-7781. 67. Awaka, J.; Kijima, N.; Hayakawa, H.; Akimoto, J. Synthesis and Structure Analysis of Tetragonal Li7La3Zr2O12 with the Garnet-Related Type Structure. J. Solid State Chem. 2009, 182, 2046-2052. 68. Wolfenstine, J.; Rangasamy, E.; Allen, J. L.; Sakamoto, J. High Conductivity of Dense Tetragonal Li7La3Zr2O12. J. Power Sources 2012, 208, 193-196. 69. Choi, J.-H.; Lee, C.-H.; Yu, J.-H.; Doh, C.-H.; Lee, S.-M. Enhancement of Ionic Conductivity of Composite Membranes for All-Solid-State Lithium Rechargeable Batteries Incorporating Tetragonal Li7La3Zr2O12 into a Polyethylene Oxide Matrix. J. Power Sources 2015, 274, 458-463. 70. Awaka, J.; Takashima, A.; Kataoka, K.; Kijima, N.; Idemoto, Y.; Akimoto, J. Crystal Structure of Fast Lithium-Ion-Conducting Cubic Li7La3Zr2O12. Chem. Lett. 2010, 40, 60-62. 71. Geiger, C. A.; Alekseev, E.; Lazic, B.; Fisch, M.; Armbruster, T.; Langner, R.; Fechtelkord, M.; Kim, N.; Pettke, T.; Weppner, W. Crystal Chemistry and Stability of “Li7La3Zr2O12” Garnet: A Fast Lithium-Ion Conductor. Inorg. Chem. 2010, 50, 10891097. 72. Rangasamy, E.; Wolfenstine, J.; Sakamoto, J. The Role of Al and Li Concentration on the Formation of Cubic Garnet Solid Electrolyte of Nominal Composition Li7La3Zr2O12. Solid State Ionics 2012, 206, 28-32. 73. Xu, B.; Duan, H.; Xia, W.; Guo, Y.; Kang, H.; Li, H.; Liu, H. Multistep Sintering to Synthesize Fast Lithium Garnets. J. Power Sources 2016, 302, 291-297.


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ACS Energy Letters

74. Xia, W.; Xu, B.; Duan, H.; Guo, Y.; Kang, H.; Li, H.; Liu, H. Ionic Conductivity and Air Stability of Al-Doped Li7La3Zr2O12 Sintered in Alumina and Pt Crucibles. ACS Appl. Mater. Interfaces 2016, 8, 5335-5342. 75. El-Shinawi, H.; Paterson, G. W.; MacLaren, D. A.; Cussen, E. J.; Corr, S. A. Low-Temperature Densification of Al-Doped Li7La3Zr2O12: A Reliable and Controllable Synthesis of Fast-Ion Conducting Garnets. J. Mater. Chem. A 2017, 5, 319-329. 76. Dumon, A.; Huang, M.; Shen, Y.; Nan, C.-W. High Li Ion Conductivity in Strontium Doped Li7La3Zr2O12 Garnet. Solid State Ionics 2013, 243, 36-41. 77. Chen, Y. T.; Jena, A.; Pang, W. K.; Peterson, V. K.; Sheu, H.-S.; Chang, H.; Liu, R.-S. Voltammetric Enhancement of Li-Ion Conduction in Al-Doped Li7-xLa3Zr2O12 Solid Electrolyte. J. Phys. Chem. C 2017, 121, 15565-15573. 78. Allen, J. L.; Wolfenstine, J.; Rangasamy, E.; Sakamoto, J. Effect of Substitution (Ta, Al, Ga) on the Conductivity of Li7La3Zr2O12. J. Power Sources 2012, 206, 315-319. 79. Song, S.; Sheptyakov, D.; Korsunsky, A. M.; Duong, H. M.; Lu, L. High Li Ion Conductivity in a Garnet-Type Solid Electrolyte Via Unusual Site Occupation of the Doping Ca Ions. Mater. Des. 2016, 93, 232-237. 80. Li, Y.; Wang, Z.; Li, C.; Cao, Y.; Guo, X. Densification and Ionic-Conduction Improvement of Lithium Garnet Solid Electrolytes by Flowing Oxygen Sintering. J. Power Sources 2014, 248, 642-646. 81. Buannic, L.; Orayech, B.; López Del Amo, J.-M.; Carrasco, J.; Katcho, N. A.; Aguesse, F.; Manalastas, W.; Zhang, W.; Kilner, J.; Llordés, A. Dual Substitution Strategy to Enhance Li+ Ionic Conductivity in Li7La3Zr2O12 Solid Electrolyte. Chem. Mater. 2017, 29, 1769-1778. 82. Gu, W.; Ezbiri, M.; Rao, R. P.; Avdeev, M.; Adams, S. Effects of Penta-and Trivalent Dopants on Structure and Conductivity of Li7La3Zr2O12. Solid State Ionics 2015, 274, 100-105. 83. Bernstein, N.; Johannes, M.; Hoang, K. Origin of the Structural Phase Transition in Li7La3Zr2O12. Phys. Rev. Lett. 2012, 109, 205702. 84. Shin, D. O.; Oh, K.; Kim, K. M.; Park, K.-Y.; Lee, B.; Lee, Y.-G.; Kang, K. Synergistic Multi-Doping Effects on the Li7La3Zr2O12 Solid Electrolyte for Fast Lithium Ion Conduction. Sci. Rep. 2015, 5, 18053. 85. Kotobuki, M.; Munakata, H.; Kanamura, K.; Sato, Y.; Yoshida, T. Compatibility of Li7La3Zr2O12 Solid Electrolyte to All-Solid-State Battery Using Li Metal Anode. J. Electrochem. Soc. 2010, 157, A1076. 86. Kim, K. H.; Iriyama, Y.; Yamamoto, K.; Kumazaki, S.; Asaka, T.; Tanabe, K.; Fisher, C. A. J.; Hirayama, T.; Murugan, R.; Ogumi, Z. Characterization of the Interface between LiCoO2 and Li7La3Zr2O12 in an All-Solid-State Rechargeable Lithium Battery. J. Power Sources 2011, 196, 764-767. 87. Ohta, S.; Kobayashi, T.; Seki, J.; Asaoka, T. Electrochemical Performance of an All-Solid-State Lithium Ion Battery with Garnet-Type Oxide Electrolyte. J. Power Sources 2012, 202, 332-335.


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88. Ohta, S.; Komagata, S.; Seki, J.; Saeki, T.; Morishita, S.; Asaoka, T. All-SolidState Lithium Ion Battery Using Garnet-Type Oxide and Li3BO3 Solid Electrolytes Fabricated by Screen-Printing. J. Power Sources 2013, 238, 53-56. 89. Du, F.; Zhao, N.; Li, Y.; Chen, C.; Liu, Z.; Guo, X. All Solid State Lithium Batteries Based on Lamellar Garnet-Type Ceramic Electrolytes. J. Power Sources 2015, 300, 24-28. 90. Jin, Y.; McGinn, P. J. Li7La3Zr2O12 Electrolyte Stability in Air and Fabrication of a Li/Li7La3Zr2O12/Cu0.1V2O5 Solid-State Battery. J. Power Sources 2013, 239, 326-331. 91. Ahn, C.-W.; Choi, J.-J.; Ryu, J.; Hahn, B.-D.; Kim, J.-W.; Yoon, W.-H.; Choi, J.H.; Lee, J.-S.; Park, D.-S. Electrochemical Properties of Li7La3Zr2O12-Based Solid State Battery. J. Power Sources 2014, 272, 554-558. 92. Yan, X.; Li, Z.; Wen, Z.; Han, W. Li/Li7La3Zr2O12/LiFePO4 All-Solid-State Battery with Ultrathin Nanoscale Solid Electrolyte. J. Phys. Chem. C 2017, 121, 14311435. 93. Thangadurai, V.; Narayanan, S.; Pinzaru, D. Garnet-Type Solid-State Fast Li Ion Conductors for Li Batteries: Critical Review. Chem. Soc. Rev. 2014, 43, 4714-4727. 94. Inaguma, Y.; Liquan, C.; Itoh, M.; Nakamura, T.; Uchida, T.; Ikuta, H.; Wakihara, M. High Ionic Conductivity in Lithium Lanthanum Titanate. Solid State Commun. 1993, 86, 689-693. 95. Belous, A. G. N., G. N.; Polyanetskaya, S. V.; Gornikov, Y. I. Study of Complex Oxides with the Composition La2/3-xLi3xTiO3. Inorg. Mater. 1987, 23, 412. 96. Bohnke, O. The Fast Lithium-Ion Conducting Oxides Li3xLa2/3−xTiO3 from Fundamentals to Application. Solid State Ionics 2008, 179, 9-15. 97. Mei, A.; Wang, X.-L.; Lan, J.-L.; Feng, Y.-C.; Geng, H.-X.; Lin, Y.-H.; Nan, C.W. Role of Amorphous Boundary Layer in Enhancing Ionic Conductivity of LithiumLanthanum-Titanate Electrolyte. Electrochim. Acta 2010, 55, 2958-2963. 98. Morata-Orrantia, A.; García-Martín, S.; Alario-Franco, M. Á. Optimization of Lithium Conductivity in La/Li Titanates. Chem. Mater. 2003, 15, 3991-3995. 99. Abreu-Sepúlveda, M.; Williams, D. E.; Huq, A.; Dhital, C.; Li, Y.; Paranthaman, M. P.; Zaghib, K.; Manivannan, A. Synthesis and Characterization of Substituted Garnet and Perovskite-Based Lithium-Ion Conducting Solid Electrolytes. Ionics 2016, 22, 317325. 100. Kwon, W. J.; Kim, H.; Jung, K.-N.; Cho, W.; Kim, S. H.; Lee, J.-W.; Park, M.-S. Enhanced Li+ Conduction in Perovskite Li3xLa2/3−x □ 1/3−2xTiO3 Solid-Electrolytes Via Microstructural Engineering. J. Mater. Chem. A 2017, 5, 6257-6262. 101. Chen, K.; Huang, M.; Shen, Y.; Lin, Y.; Nan, C. Improving Ionic Conductivity of Li0.35La0.55TiO3 Ceramics by Introducing Li7La3Zr2O12 Sol into the Precursor Powder. Solid State Ionics 2013, 235, 8-13.


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102. Kotobuki, M.; Suzuki, Y.; Munakata, H.; Kanamura, K.; Sato, Y.; Yamamoto, K.; Yoshida, T. Fabrication of Three-Dimensional Battery Using Ceramic Electrolyte with Honeycomb Structure by Sol-Gel Process. J. Electrochem. Soc. 2010, 157, A493-A498. 103. Li, C.-L.; Zhang, B.; Fu, Z.-W. Physical and Electrochemical Characterization of Amorphous Lithium Lanthanum Titanate Solid Electrolyte Thin-Film Fabricated by EBeam Evaporation. Thin solid films 2006, 515, 1886-1892. 104. Le, H. T.; Ngo, D. T.; Kim, Y.-J.; Park, C.-N.; Park, C.-J. A Perovskite-Structured Aluminium-Substituted Lithium Lanthanum Titanate as a Potential Artificial SolidElectrolyte Interface for Aqueous Rechargeable Lithium-Metal-Based Batteries. Electrochim. Acta 2017, 248, 232-242. 105. Stramare, S.; Thangadurai, V.; Weppner, W. Lithium Lanthanum Titanates: A Review. Chem. Mater. 2003, 15, 3974-3990. 106. Bruce, P. G.; West, A. The A‐C Conductivity of Polycrystalline LISICON, Li2+2xZn1-xGeO4, and a Model for Intergranular Constriction Resistances. J. Electrochem. Soc. 1983, 130, 662-669. 107. Robertson, A.; West, A.; Ritchie, A. Review of Crystalline Lithium-Ion Conductors Suitable for High Temperature Battery Applications. Solid State Ionics 1997, 104, 1-11. 108. Thangadurai, V.; Weppner, W. Recent Progress in Solid Oxide and Lithium Ion Conducting Electrolytes Research. Ionics 2006, 12, 81-92. 109. Kuwano, J.; West, A. R. New Li+ Ion Conductors in the System, Li4GeO4Li3VO4. Mater. Res. Bull. 1980, 15, 1661-1667. 110. Deng, Y.; Eames, C.; Fleutot, B.; David, R. n.; Chotard, J.-N. l.; Suard, E.; Masquelier, C.; Islam, M. S. Enhancing the Lithium Ion Conductivity in Lithium Superionic Conductor (LISICON) Solid Electrolytes through a Mixed Polyanion Effect. ACS Appl. Mater. Interfaces 2017, 9, 7050-7058. 111. Song, S.; Lu, J.; Zheng, F.; Duong, H. M.; Lu, L. A Facile Strategy to Achieve High Conduction and Excellent Chemical Stability of Lithium Solid Electrolytes. RSC Adv. 2015, 5, 6588-6594. 112. Kato, T.; Hamanaka, T.; Yamamoto, K.; Hirayama, T.; Sagane, F.; Motoyama, M.; Iriyama, Y. In-Situ Li7La3Zr2O12/LiCoO2 Interface Modification for Advanced AllSolid-State Battery. J. Power Sources 2014, 260, 292-298. 113. Zhou, W.; Wang, S.; Li, Y.; Xin, S.; Manthiram, A.; Goodenough, J. B. Plating a Dendrite-Free Lithium Anode with a Polymer/Ceramic/Polymer Sandwich Electrolyte. J. Am. Chem. Soc. 2016, 138, 9385-9388. 114. Li, Y.; Xu, B.; Xu, H.; Duan, H.; Lü, X.; Xin, S.; Zhou, W.; Xue, L.; Fu, G.; Manthiram, A. Hybrid Polymer/Garnet Electrolyte with a Small Interfacial Resistance for Lithium‐Ion Batteries. Angew. Chem., Int. Ed. Engl. 2017, 129, 771-774.

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1) The additional Li+ are introduced to Li1.2Al0.2Ti1.8(PO4)3 by partially substituting Ti4+ with Al3+, leading to structural changes and influencing the Li+ mobility 2) The increase in pH of LLZO powder/water suspension was resulted from H+/Li+ exchange reaction between water and the LLZO electrolyte. 3) Though perovskite-type conductors have high bulk conductivities, the total conductivities were lower (~2 × 10−5 S cm−1 at room temperature) due to higher grain boundary resistances 4) Insertion of polymer layer between SE and anode suppresses dendrite formation due to uniform Li+ flux across the polymer/lithium interface


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