Interface in Solid-State Lithium Battery: Challenges, Progress, and

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Cite This: ACS Appl. Mater. Interfaces 2019, 11, 22029−22050

Interface in Solid-State Lithium Battery: Challenges, Progress, and Outlook Syed Atif Pervez,*,† Musa Ali Cambaz,† Venkataraman Thangadurai,‡ and Maximilian Fichtner† †

Helmholtz Institute Ulm, Helmholtzstraße, 11, Ulm89081, Germany Department of Chemistry, University of Calgary, 2500 University Drive Northwest, Calgary, Alberta T2N 1N4, Canada

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‡

ABSTRACT: All-solid-state batteries (ASSBs) based on inorganic solid electrolytes promise improved safety, higher energy density, longer cycle life, and lower cost than conventional Li-ion batteries. However, their practical application is hampered by the high resistance arising at the solid−solid electrode−electrolyte interface. Although the exact mechanism of this interface resistance has not been fully understood, various chemical, electrochemical, and chemomechanical processes govern the charge transfer phenomenon at the interface. This paper reports the interfacial behavior of the lithium and the cathode in oxide and sulfide inorganic solid-electrolytes and how that affects the overall battery performance. An overview of the recent reports dealing with high resistance at the anodic and cathodic interfaces is presented and the scientific and engineering aspects of the approaches adopted to solve the issue are summarized. KEYWORDS: interface resistance, solid-electrolyte, Li dendrites, solid-state battery, interfacial characterization

1. INTRODUCTION The Paris agreement in 2015 signifies the world’s efforts in combating the issue of global warming by targeting to keep the increase in global average temperature below 2 °C.1 To achieve this goal, development of technologies with ideally no or reduced greenhouse gas emission is needed for gridelectrification and automotive transportation. Renewable energy storage using Li-ion battery (LIB) is a viable option to achieve the aforementioned goal. In the past decade, the research in LIB has gained significant momentum for practical applications in portable electronics market and electric vehicle (EV) industry.2,3 Conventional LIBs consist of intercalation cathodes (e.g., LiCoO2, LiFePO4), anode (graphitic carbon) and organic solvent-based liquid electrolytes. The energy density of batteries equals the product of the working potential, which is the differential between anode and cathode potential and the specific capacity. In order to increase the energy density either the anode materials should operate at the lowest possible potentials (0 V vs Li/Li+) and/or cathode materials should operate at the highest possible potentials (≤5 V vs Li/ Li+).4 The liquid electrolytes provide an ionic conductivity in the order of 10−2 S cm−1, which is sufficient to shuttle the ions between the electrodes during charge/discharge cycles. The conventional LIBs may continue to fulfill the energy demands in stationary storage systems; however, their role in highenergy applications such as EVs and portable electronics is limited because of various issues with liquid electrolytes such as their limited potential window, lower Li ion transference number, and flammable nature.2,5,6 Furthermore, utilization of Li metal poses serious risk in liquid systems owing to their instability with organic electrolytes.5,7,8 It is widely observed that during repeated charge/discharge cycles, Li metal forms © 2019 American Chemical Society

mossy or needle-shaped structures, possibly due to varying electric fields originating from irregular surface morphologies, eventually short circuiting the battery. These issues can be addressed by replacing liquid electrolytes with solid-state electrolytes (SSEs). SSEs are considered safer due to their nonflammable nature.9 This reduces safety features on pack level, which can enhance the overall energy density of the system. Further, suppression of Li dendrites may be possible owing to the high shear modulus of such electrolytes.10,11Among various types of SSEs including solid polymers and polymer−ceramic composites, inorganic SSEs have drawn an enormous attention, of late. There are various structural families in inorganic SSEs which can be broadly divided in to oxides, sulfides, and argyrodites.12−14 Oxides include Li− garnets,15,16 perovskites,17,18 Li super ionic conductors (LISICONs),19−21 and Na super ionic conductors (NASICONs).2,22 Sulfides include thio-LISICON-type LGPS23,24 and glassy-type Li2S−P2S5.25−27 Argyrodites exhibit chemical composition of Li6PS5X (X = Cl, Br, I).28 When compared with polymer SSEs, inorganic SSEs offer higher ionic conductivities, better thermal stabilities and much higher Li transference number (tLi+∼ 1 vs 0.2−0.5 for polymers). However, despite the potential advantages, realization of a solid-state battery based on inorganic SSEs with comparable electrochemical performance to conventional LIBs is not achieved. This is mainly due to their relatively lower ionic conductivities than liquid electrolytes and high solid−solid interface resistance between electrodes and SSEs. In recent Received: February 14, 2019 Accepted: May 30, 2019 Published: May 30, 2019 22029

DOI: 10.1021/acsami.9b02675 ACS Appl. Mater. Interfaces 2019, 11, 22029−22050

Review

ACS Applied Materials & Interfaces

inorganic SSEs work on the same principle as their liquid counterparts. During a charge cycle, Li ions eject from the cathode lattice structure, move through the inorganic SSE, and finally plate onto the Li. Over the years, numerous inorganic SSEs were investigated.12,14 Initially, the efforts were mostly focused on increasing the ionic conductivities by structural and compositional tuning of the materials. In general, for an inorganic SSE to have high ionic conductivity, it should possess the following:29 • a large number of empty lattice sites either through vacancies or interstices to ensure high mobility of the ions; • low migration enthalpy to promote jumping and hopping of mobile ions; • preferably a 3D solid framework with channels allowing “molten” migrating ions without structure destruction; • anion with high polarizability to accommodate migrating ions through covalent bonds. The ion transport in inorganic SSEs is quantified by Arrhenius equation, i.e.,

years, a significant progress has been made in achieving reasonably high ionic conductivities in inorganic SSEs. However, the issue of high interface resistance is not solved yet. Although it is difficult to prove experimentally the exact mechanism for interface resistance, current studies indicate that chemical incompatibility, electrochemical instability, and mechanical issues between electrodes and SSEs may be responsible for such interfacial phenomenon (Figure 1).

i A y i −E y σ = jjj zzzexpjjj a zzz k T { k kT {

(1) −1

where σ is the conductivity of the electrolyte in S cm , A the frequency factor, Ea activation energy expressed in joules per mole, k Boltzmann’s constant, and T temperature measured in Kelvin. The integration of the SSE with various electrodes should be chemically stable at a wide range of potentials to build a working battery.30 Further, the electrolyte itself should be cost-effective, thermally stable, and mechanically robust. It is very difficult for a single type of solid electrolyte to offer all the aforementioned advantages. Table 1 lists some of the known SSEs and their merits and demerits. Garnet Type. Parent garnet-type Li ion conductors with the general chemical formula Li5La3M2O12 (M = Nb, Ta) were initially developed by Thangadurai et al. with a roomtemperature ionic conductivity of 1 × 10−6 S cm−1.31 Zr and Ta-based Li−stuffed garnets demonstrated good chemical stability at various temperatures. The conductivity was further increased (1 × 10−3 S cm−1 at room temperature) by realizing lower activation energies for Li-ion hopping (0.35−0.4 eV) through substituting La and M sites with various metal ions. Lirich garnets were obtained such as Li6ALa2M2O12 (A = Mg, Ca, Sr, Ba)32 and Li7La3C2O12 (C = Zr, Sn).33 The improved ionic conductivities in Li-rich garnets were attributed to the presence of excess Li ions in distorted octahedral sites (48g/ 96h) rather than confining only in the tetrahedral sites (24d).34 Li garnets offer various advantages such as high thermodynamic stability vs Li/Li+ and better mechanical strength and thermal stability at higher temperatures. However, the shortcomings are the formation of insulating carbonates and hydroxide layers on the surface of Li garnets when exposed to air and humid environments and their brittle nature, which hampers device integration. Perovskites. Perovskites have the general formula ABO3 (A = Ca, Sr or La; B = Al, Ti).17 By aliovalent doping at A site, Li can be introduced in the structure resulting in perovskitetype Li3xLa(2/3−x)□(1/3−2x)TiO3 (0 < x < 0.16) (LLTO).35,36 These compounds have shown a high bulk ionic conductivity up to 1 ×10−3 S cm−1 at room temperature with activation energy of ∼0.3 eV. Highly conductive Ti-based perovskite SSEs are chemically stable in air/humid conditions and in a

Figure 1. Interface in solid-state battery. “Magnified views” are schematic illustration of the Li and cathode interface regions. Issues with each interface and approaches toward solving them are listed.

In this review, we begin with a general introduction of ASSB, followed by short descriptions of various types of inorganic SSEs. In the following section, the mechanism of formation of interphase layers due to (electro-)chemical redox reactions taking place at the electrode-SSE junction is explained with the help of an energy band diagram. Further, the mechanical issues due to the stresses arising at the interfaces are also discussed. We focus on anodic and cathodic interface in oxide and sulfide SSE systems, and survey recent reports investigating the factors behind the formation of interphase layers and outlining strategies to tackle the issues. We review recent works, which investigate the role of high Li-SSE interface resistance in initiating Li dendrites. Further hybrid liquid−solid systems for advanced batteries are discussed. We also review in situ and operando characterization techniques currently used to study the dynamic processes happening at the electrode−SSE interfaces.

2. BRIEF OVERVIEW OF INORGANIC SOLID ELECTROLYTES The inorganic SSEs are mostly ceramic Li-ion conductors with negligible electronic conductivity and hence they can act as separator between the electrodes. Solid-state batteries based on 22030


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Table 1. Summary of the Crystal Structures, Ionic Conductivity, Activation Energy, Performance, and Limitations of Various Inorganic Solid Electrolytes

Further, their synthesis involves high-temperature sintering to decrease their high grain-boundary resistance.35 LISICON Type. Lithium super ionic conductors (LISICONs) are based on crystal framework similar to the γ-Li3PO4 structure with an orthorhombic unit cell and Pnma space

wide range of temperatures. However, common concerns associated with these materials are that they undergo facile reduction vs Li at voltages ≤1.5 V, which makes them unsuitable with negative electrodes operating at low voltages. 22031


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Figure 2. Schematic illustration of open-circuit energy diagram for a Li-SSE-LixMyO2 solid-state battery system. μLi and μLixMyO2 represent the chemical potentials of Li and the cathode materials, respectively. ΦLi is the electrostatic potential of Li and ΦLixMyO2 is the potential of LixMyO2.2

NASICON Type. Na super ionic conductors (NASICONs) are solid-state ionic conductors with general formula LiM2(XO4)3(M = Ge, Ti, or Zr; X = S, P, As, Mo). Generally the compounds are rhombohedral structures (space group R¯3c) with a 3D framework built up by M2(XO4)3 units in which two MO6 octahedra and three XO4 tetrahedra share oxygen atoms while the A+ ions diffuse through interstices.2,45,46 In titanium-based NASICON-type materials, the Li+ ion conductivity is greatly enhanced when Ti4+ is partially replaced by trivalent cations (Al, Ga, Sc, In, Y).22,47 The maximum ionic conductivity for NASICON type obtained are ∼1 × 10−3 S cm−1 at 25 °C for Li1+xTi2−xAlx(PO4)3 (LATP) where x ≈ 0.3. High ionic conductivity makes NASICONs an attractive choice. However, when in contact with Li metal, they undergo reduction reaction (Ti4+ to Ti3+,) which is a major drawback.48 Li2S−P2S5 Type Glassy Sulfide. The earliest glassy type sulfide solid-electrolytes were Li2S-SiS2 system.49 Various other types which followed are LiS-GeS2, Li2S−B2S3 and Li2S−P2S5. Among them the quasi-binary system xLi2S(100-x)-P2S5(x from 70 to 80) is of particular interest because of high ionic conductivity (10−3−10−4 S cm−1).50,51 Such sulfide glassceramic SSE family has a 3D framework structure and 1D Li conduction path along the c-axis.52 Despite their high ionic conductivity, a major disadvantage is that they tend to generate poisonous H2S gas when exposed to ambient air.53 Argyrodite. The general chemical formula of Li argyrodites based on halogens is Li6PCh5X where Ch = O, S or Se and X = Cl, Br or I.28,54,55 The structure is based on tetrahedral close packing of anions (crystallizes in cubic F43m space group) where P atoms coordinate with S to form PS4 tetrahedrons while Li ions are distributed over the tetrahedral interstices (48h and 24g sites).56 DFT calculations showed that higher

group. Early reports from Bruce and Roberston showed feasibility of Li2+2xZn1−xGeO4 as solid-state Li conductors.37,38 LISICONs are three-dimensional structures with a hexagonal close packing of oxygen atoms with Ge, Li and Zn cations occupying the tetrahedral and octahedral interstices.39 Major advantages of LISICON-type electrolytes are their thermal stability and compatibility with aqueous electrodes.40 However, the room-temperature conductivity of these compounds is quite low (∼1 × 10−7 S cm−1). Further, thermodynamic instability with Li and reactivity with CO2 are major concerns with such types of electrolytes.14 Thio-LISICON. In thio-LISICON SSEs, the O2− in LISICON-type is replaced by bigger and more polarizable S2− anion which improves the conductivity significantly.41 Li10MP2S12 (M = Si, Ge, Sn) is commonly known as LGPS thio-LISICON solid-state ion conductors reported by Kamaya et al. with a breakthrough conductivity of 1.2 × 10−2 S cm−1 at 25 °C.23 This superionic conductor has a 3D framework with tetragonal unit cell made of PS4 and GeS4 tetrahedra structure and space group P42/nmc. The Li ion conduction is both onedimensional along the c-axis and two-dimensional in the abplane with 0.16 and 0.26 eV activation energies, respectively.42 These sulfides offer high ionic conductivity in comparison to oxides because of low grain-boundary resistance. Also, they are mechanically robust because of their ductile nature. However, they are costly because of presence of Ge; therefore, other metals such as Sn and Si are employed to reduce the cost.43 LGPS electrolytes are not stable against reduction of Li metal, where they decompose to form insulating products (Li3P, Li2S, Li15Ge4) that increase the interface resistance.13 Also, they are highly moisture sensitive and possibly hazardous because of the formation of H2S gas if exposed to air for long periods of time.44 22032


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resulting in continuous decomposition of the SSE. The decomposed products will form thicker and electrochemically unstable interphase layers, which will block/hinder ionic transport at the interface, eventually resulting in higher resistance. Such is the case with sulfide SSEs when integrated with Li and metal oxide cathodes.65−68 Ge-based sulfide SSEs such as LGPS are unstable with Li because they start reducing at ∼1.7 V and oxidizing at ∼2.15 V, leaving a very small potential window for thermodynamically stable operation of the battery. Here, formation of electronically conductive layers has been observed at the interface, which continuously favors the decomposition reaction of the SSE.67 Similarly, with LiCoO2 (LCO) cathodes in a sulfide SSE system, mutual diffusion of S and Co species will occur at the interface resulting in electronically conductive cobalt sulfide layers.68 Hence, such layers both at the Li and at the cathode side cannot provide the required passivation, which is very critical for a stable interface. On the other hand, formation of an electron-insulating interphase layer will passivate the SSE and hence inhibits further decomposition. A prerequisite for such phenomenon to occur is that VOC should not be too high compared to the potential window of the electrolyte.2 With the help of electronically insulating interphase layers, the higher potentials of the electrode can be adjusted to the potential limits of the SSE (Figure 2) that will provide kinetic stability to the ion transport at the interface. Barring a few oxide SSEs,69,70 most show better electrochemical compatibility with Li and oxide cathodes than sulfide SSEs. The reduction and oxidation potentials for oxide SSEs are generally between 0.5 and 3 V, which provides wider potential window than sulfide SSEs. Among known SSEs, Li7La3Zr2O12 (LLZO) has arguably the best stability with Li. Unlike sulfide SSEs, they undergo reduction reaction at much lower potentials (∼0.05 V vs Li/ Li+). At such potential, the reaction products (Li2O, Zr3O, and La2O3) may passivate the SSE, which helps in forming interphase layers with optimum thickness.13 It is important to mention that the formation of such passivation layers is strongly dependent on the interfacial surface chemistry and the type of dopant in LLZO. A recent study demonstrated that LLZO with surface carbonates could not be wet with Li. Furthermore, unlike Ta and Al-doped LLZO samples, Nbdoped LLZO undergo continuous reduction reaction upon contact with Li, hence proper passivation is not achieved which results in continuous increase in interface resistance.71 On the cathode side, few theoretical studies have demonstrated good stability of LLZO with cathodes, in particular LiCoO2.13,72 Ceder et al. carried out DFT calculations to conclude that among various high voltage cathode materials (e.g., LiCoO2, LiNi0.5Mn1.5O4, and LiFePO4), LCO shows the most stable behaviors with LLZO. Experimentally, however, there are contradicting claims regarding the stability of metal-oxide cathodes with LLZO. Kotobuki et al.73 suggested good compatibility between LCO and Ta-doped LLZO. Few other works74,75 reported formation of undesirable interphase layers. Kim et al. observed that the Li-ion transport suffered because of the formation of a La2CoO4 based interphase layer at the LCO-LLZO junction. Further, in the case of LiNi0.5Mn1.5O4 (LNMO) cathode materials and LLZO, Hansel et al. observed a potential drop at ∼3.8 V during the first charge cycle. This was due to formation of inactive phases at the LNMO−LLZO interface that eventually led to the failure of the cell.

ionic conductivity could be achieved by increasing the halogens content in the compounds and optimizing their distributions in the crystal lattice.57 Li argyrodites, especially Li6PS5Cl and Li6PS5Br compounds, have shown Li ionic conductivities greater than 10−3 S cm−1.56,58,59 Further they offer lower Ea (0.16 eV-0.56 eV) and lower synthesis cost. However, the electrolyte decomposition has been observed when in contact with Li metal. The decomposed products (Li3P, Li2S and LiX) formed a continuously growing interphase layer, with a corresponding increase in interfacial resistance.60 Further, as with other sulfide type electrolytes, their sensitivity to humid environments is a major concern.

3. INTERFACIAL PHENOMENA AT ELECTRODE SOLID−ELECTROLYTE INTERFACE Understanding the mechanism behind the interfacial processes in solid-state batteries is challenging because of the complex nature of the interface between electrode active materials and the SSEs. The origin of the interfacial phenomenon in a solid system is determined by two important factors, first, the chemical and electrochemical compatibility between the electrodes and SSEs, and second, the mechanical robustness of the interface contact. Chemical reactions may occur at open circuit potentials or under biased conditions and result in the formation of interphase layers at the electrode−electrolyte junctions.2,13 In general, a SSE is considered to be stable when it undergoes electrochemical reactions with electrodes to an extent that passivating interphase layers with finite thickness are achieved. On the other hand, mechanical stability of the interface is also important because fractures could result in loss of electrode−electrolyte contact, which will create a barrier for Li transport.61,62 Such mechanical stresses are expected because of the continuous change in lattice of electrodes during Li (de)insertion, which may delaminate the electrodes from the electrolyte. A schematic of the open circuit energy band diagram of a solid-state system based on Li metal anode, solid electrolyte, and metal oxide cathode (LixMyO2) is presented in Figure 2. The electrochemical potential window is defined by the difference of the reduction and oxidation potentials of the solid electrolyte. Voltages applied beyond this window will reduce and oxidize the electrolyte on the Li side and the LixMyO2 side, respectively. μLi and μLixMyO2 represent the chemical potentials of Li and the cathode materials, respectively. ΦLi is the electrostatic potential of Li anode. ΦLixMyO2 is the potential of the metal-oxide cathode. The chemical potential difference between Li and LixMyO2 represents the open circuit voltage (VOC) of the battery, i.e., VOC = μLi − μLi M O x

y 2

(2)

Ideally, the potentials of the electrodes should be within the potential window of the electrolyte to achieve stable electrochemical performance.2,63,64 However, in a scenario as shown in Figure 2, where μLi and μLixMyO2 are beyond the potential window of the electrolyte, formation of anode and cathode interphase layers will take place because of the decomposition of the SSE. Depending on the type of the SSE and difference between VOC and the potential window of the electrolyte, the interphase layers may either be electronically conductive or insulating.13 In case the interphase layer permits electronic conductivity, passivation will not be achieved, 22033


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Figure 3. (a, b) SEM image and elemental mapping of ZnO-coated LLZO surface. (c) Illustration of Li interaction with ZnO-coated garnet LLZO. (d, e) Digital photos of the front and back side of garnet pellet after lithiation. (f) SEM image of the Li−garnet junction showing the improved interface contact. Panels a−f reprinted with permission from ref 81. Copyright 2016 American Chemical Society. (g) Schematic representation of Li−Mg alloy formation. (h) SEM images showing dissolution of Mg in Li metal forming Li−Mg alloys at the Li-garnet interface. (i) Li stripping/ plating voltage profiles at a current density of 0.1 mA cm−2. Panels g−i reprinted with permission from ref 83. Copyright 2016 Wiley.

4. LITHIUM ANODE−SSE INTERFACE 4.1. Improving Li−SSE Contact with Thin Coatings. The poor physical contact due to the brittle and rigid nature of oxide SSEs with Li, in particular LLZO, leads to high interface resistance.78 Pressing Li at high pressures or melting it on top of the LLZO pellets has helped in improving the electrode− electrolyte contact.79,80 A better strategy is coating the LLZO surface with a thin layer (thickness: 10−20 nm) of materials having good chemical reactivity with Li. Because of the reaction between Li and the coated materials, the LLZO surface could be “wetted” efficiently, resulting in improved contacts. Various materials, which could form alloys with Li, were coated on LLZO SSEs through high vacuum techniques such as sputtering, atomic layer deposition (ALD), pulse enhanced chemical vapor deposition (PECVD), and chemical vapor deposition (CVD). For example, coating LLZO SSE with a few-nanometer-thick ZnO layer through ALD, significantly improved its wettability.81 As evident from Figure 3a, b, a 30−50 nm thick ZnO layer is deposited on the surface of LLZO. On contact with Li, the ZnO is reduced to Zn metal and forms Li−Zn alloys. During the alloying process, Li

In terms of mechanical properties, sulfide SSEs have a slight edge over oxide SSEs thanks to their ductile nature that helps in their integration into bulk batteries by simple technique such as cold pressing. Nevertheless, various studies62,76,77 have shown that loss of contact still occurs between active materials and the sulfide SSE during cycling. For example, in a LiNi0.8Co0.1Mn0.1O2 solid-state battery with β-Li3PS4, a severe capacity loss was observed that was partly blamed to the decomposition of the SSE during first cycle and partly to the electrode−electrolyte contact loss due to chemo-mechanical contraction of the active material upon delithation.76 Quantitative analysis of Li2S−P2S5 by an electro-chemomechanical model confirmed the loss of contact due to mechanical fracture. It was suggested that SSEs with Young’s modulus ≥15 GPa are more likely to get microcracking. Understandably, this issue gets more severe in the stiff and brittle oxide SSEs with much higher Young’s modulus (LATP, LLTO, and LLZO have approximately 115, 193, and 150 GPa, respectively) than sulfide-type.10 22034


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Figure 4. (a) 3D architecture printed through 3D printing technique on LLZO. (b) Schematic and SEM image showing Li-filled pores on LLZO substrate; Li stripping/plating voltage profiles of the SSE. Panels a, b reprinted with permission from ref 92. Copyright 2018 Wiley. (c) SEM images showing porous-dense-porous layers of garnet SSE; schematic figures of the mechanism of Li plating on Cu substrate. Reprinted with permission.93 Copyright 2018, National Academy of Sciences. (d) Contact angle measurements of molten Li on polished and nonpolished LLZO garnet pellets. Reprinted with permission.94 Copyright 2017 American Chemical Society.

LiH2PO4-coated Li. Lower interface resistance translated into much improved performance of the full cell, which demonstrated initial discharge capacity of 131 mAh g−1 with a capacity retention of ∼85% after 500 cycles at a rate of C/10. 4.2. Surface Modifications of the SSE. Wachsman et al. printed 3D architectures on LLZO surface using 3D printing technique (Figure 4a, b).92 The work was aimed at innovative architectural designing of the surface with a micrometer-scale 3D pattern, which could increase areal contact between Li and SSE. To demonstrate feasibility of the structures, Li-symmetrical cells were fabricated (Figure 4b) and cycled reversibly for 150 h at current density of 0.33 mA cm−2. A similar concept is presented in another work, where a porous-denseporous garnet structure was fabricated in a layer-by-layer design (Figure 4c)93 It was proposed that the porous structures could act as 3D Li-ion conductive framework for Li and sulfur electrodes. The feasibility of the structure was demonstrated by infiltrating Li in the upper porous layer, the dense middle layer acted a separator while the bottom empty layer was gradually filled with Li by electrochemical plating (Figure 4c). The Li symmetrical cells exhibited cycling stability with a capacity of 1 mAh cm−2 at 0.5 mA cm−2 for 300 h without any dendrite formation or significant overpotential. Lei Cheng and co-workers showed that surface microstructure of the SSE with smaller grain size (20−40 μm) offers lower interfacial resistance (37 Ω cm−2) that helps in achieving better cycling performance.79 In another approach, LLZO surface was “conditioned” by various polishing techniques to get rid of

diffuses through the metal-oxide layer and improves the Li− LLZO contact. Schematic Figure 3c shows the process of lithiation of the ZnO layer, also shown in photos (Figure 3d, e). The SEM image in Figure 3f shows no voids at the Li− LLZO pellet interface, suggesting improved electrode−electrolyte contact, which was supported by EIS analysis (∼ 20 Ω cm−2). Recently, Zhou et al. coated ZnO on LLZO by a wet chemical surfactant method, which is easily scalable, unlike vacuum deposition methods, and observed an interface resistance of 10 Ω cm−2 at room temperature.82 It was proposed that coating a thin Mg layer at the interface between Li and LLZO helps in suppressing the interface resistance.83 By heating the Li−Mg-LLZO interface, the Mg diffuses in to the bulk Li forming Li−Mg alloys that improved the Li-wettability (Figure 3g, h). The Li-symmetrical cells showed much lower interface resistance and stable Li stripping/plating voltage profiles for an extended period of time (Figure 3i). Similar strategies have been adopted to form Li alloys at the interface by coating Si,78,84 Au,85 Al2O3,86 Al,87 Ge,88 and graphite.89 Nagao et al. stabilized the interface between Li and Li2S−P2S5 SSE by forming a thin indium (In) layer on the surface of the SSE by vacuum-evaporation.90 The Li−In alloy offered improved electronic conductivity at the interface and reduced the overpotential and ensured low interface resistance. In another work, a LiH2PO4 layer was formed on the surface of Li. The layer served as protection for LGPS SSE, which otherwise would have reduced upon contact with bare Li.91 The EIS analysis showed much less interface resistance for 22035


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Figure 5. (a) Schematic representation of the Li-SPE-LAGP-LFP cell. (b) SEM image showing cross-section of the cell. (c) Cyclic performance and Columbic efficiency of the cell. Panels a−c reprinted with permission from ref 102. Copyright 2017 Royal Society of Chemistry. (d) Schematic representation of the gel-modified cell structure. (e) Battery performance of the Li-gel-LLZO-gel-LFP cell. Panels d, e reprinted with permission from ref 104. Copyright 2017 American Chemical Society.

0.1 C at 60 °C was achieved (Figure 5c). Despite the aforementioned advantages of PEO polymer, one of the major shortcomings is their low Li+ conductivity (∼1 × 10−6 S cm−1 at 25 °C). Therefore, these polymers are operated at ≥60 °C to achieve high ionic conductivities. Goodenough et al. argued that at such high temperatures the PEO polymer does not remain in the film form but rather transforms into a gel that could diffuse across the interphase layer, leading to loss of Li+ from cathode.103 In addition, because of the low transference number (0.2−0.5) of PEO polymers, rapid anion depletion can occur at the interface, building a large electric field that may favor decomposition of the electrolyte and enhanced Li dendrite nucleation. To address the issues, they used a crosslinked poly(ethylene glycol) methyl ether acrylate (CPMEA) polymer, thermally stable at much higher temperatures (270 °C) and ionically conductive (∼1 × 10−4 S cm−1 at 65 °C), which showed good specific capacity and stable cyclic performance for more than 600 cycles. The improved battery performance was attributed to the polymeric interlayers, which helped in regulating the electric field across the interface. PVDF-HFP based gel electrolyte was used as an interlayer between Li-LLZO and LLZO-LFP interfaces (Figure 5d).104 They took advantage of the soft nature, high ionic conductivity, and better wettability offered by the polymer to reduce the anodic and cathodic interfacial resistance with the SSE to 248 and 214 Ω cm2, respectively. Further the cell demonstrated a specific capacity of ∼140 mAh g−1 for more than 70 cycles at room temperature and Coulombic efficiency of ∼ 90% (Figure 5e). 4.4. Li Dendrite Initiation Due to Poor Li−SSE Interface. Considering the Monroe and Newman model for polymer SSEs,105 where they proposed a minimum limit on the shear modulus (≥6 GPa) for blockage of Li dendrites, ceramic SSEs should be more efficient in suppressing Li dendrite formations because of their higher shear modulus than polymer SSEs. However, various studies79,84,85,106−111 have shown formation of Li dendrites in crystalline SSEs, which eventually

Li2CO3 and LiOH that can grow on the surface upon exposure to air and humidity.94−96 By polishing off the Li2CO3 surface layer, wettability of the SSE was improved as evident from a lower contact angle between molten Li and LLZO (Figure 4d).94 Goodenough’s group argued that insulating carbonate and hydroxide layers grow not only on the surface of the Li garnet SSE but also on the individual grains in bulk of the material. Therefore, they incorporated LiF during solid-state synthesis of Li garnet materials. LiF has a strong ionic bond and is insoluble in water, hence it can resist the formation of carbonates and hydroxides in humid air not only on the surface but also on the individual grains of the SSE.97 In another work, LLZO surface was cleaned from carbonate contaminations by making carbon react with Li2CO3 at 700 °C.98 Through mass spectrometric analyses, the authors observed evolution of CO in the garnet/carbon mixture at ∼700 °C. They argued that carbon reacts with Li2CO3 on the garnet to form CO and Li2O (Li2CO3 + C → Li2O + 2CO). Overall, strategies to physically modify the surface of the SSEs are considered simple and costeffective in contrast to vacuum-based coating techniques. 4.3. Polymeric Interlayer at the Interface. Because of their flexible and cohesive nature, Li-ion conducting polymer layers at the electrode−SSE interface are a good choice to avoid the solid−solid contact between electrode and SSE. Poly(ethylene oxide) (PEO) generally exhibits good electrochemical stability with Li and excellent compatibility with Li salts.99,100 Inspired by the positive role of PEO film in stabilizing the Li−electrolyte interface and suppression of Li dendrite in liquid systems,101 Yang et al. adopted a similar strategy in solid-state systems, where they coated a NASICONtype SSE by PEO-based solid polymer electrolyte (SPE).102 The cell comprising Li, SPE, SSE pellet and LFP cathode was assembled in a multilayer form (Figure 5a, b). The PEO interlayer significantly reduced the interface resistance, which resulted in stable Li stripping/plating voltage profiles with much lower voltage polarization. Further, at a high reversible capacity of 150 mAh g−1, stability for more than 200 cycles at 22036


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Figure 6. (a) Schematic showing variations in current and potential contours due to interface resistance; current density profile highlighting abrupt increase in current density due to roughness on the Li surface. Reprinted with permission from ref 108. Copyright 2016 Elsevier. (b) Schematic illustration of Li-SSE contacts and regulation of Li flux due to Au buffer layers at the interfaces. Reprinted with permission from ref 85. Copyright 2017 American Chemical Society. (c) Schematic figure highlighting the role of Li3PO4 in suppressing Li dendrites in Li-LLZO cells. Reprinted with permission from ref 112. Copyright 2017 Elsevier.

conducting interphase layer will help in alleviating Li dendrite growth at the interface. They justified their hypothesis by adding a small amount of liquid electrolyte at the Li−SSE interface. The liquid electrolyte was able to form a passivating SEI. Consequently, a significant retardation in the formation of Li dendrites was observed. They also modified the LLZO surface by filling the microstructures with Si nanoparticles by a facial polishing process. Li-ion reaction with Si formed Li−Si alloys that stabilized the interface and helped in retarding Li dendrites. Raj and co-workers presented the importance of ionic conductivity and fracture resistance of the solid electrolyte and current density on the formation of Li dendrites.108 It was shown that higher grain-boundary resistance of the electrolyte and closer proximity of the grain boundaries to the Li interface will promote Li dendrite formation. The model also emphasized the role of the Li−SSE interface on nucleation of Li dendrites and demonstrated that a roughness on the Li surface at the interface will enhance the local electric field that could provide nucleation sites, favoring further formation of Li dendrites. The roughness on the Li surface can be easily formed because of poor contact, where Li is interfaced with the SSE only at limited positions, which act as hotspots due to large concentration of Li ions. As shown in Figure 6a, a bump on Li surface will concentrate the current density enhancing the electric field at the tip. Depending on the aspect ratio (width w and height h) of the bump, a much higher current density (j(h/w)) builds up, which is responsible for the dendrite formations. Assuming a 100 MPa fracture stress, bump aspect ratio of unity, and current density of 1 mA cm−2,

lead to short circuiting despite their relatively high shear modulus than polymer and liquid electrolytes. Porz et al.110 conducted a study on amorphous Li2S−P2S5, polycrystalline βLi3PS4, LLZO, and single crystalline LLZO to understand the mechanism of Li dendrite growth and propagation. They did not observe any dendrite growth in amorphous Li2S−P2S5, although its presence was confirmed in crystalline SSEs. They claimed that Li dendrites propagate through the defects, pores, and cracks in the crystalline SSEs, regardless of their shear modulus. Han et al. investigated the isolated formations of Li dendrites in SSEs and suggested that the electronic conductivity in LLZO and Li2S−P2S5 solid electrolytes is responsible for the nucleation and growth of Li directly in bulk of the samples.111 On the basis of their experimental results, they proposed a mechanism for the growth of Li dendrites, which was based on two critical conditions, first, the presence of mobile electrons inside the electrolyte, and second, the potential difference between the bulk of electrolyte and Li electrode. In an electronically conductive solid electrolyte, the electrons will combine with Li+ to deposit Li and the presence of these electrons will decrease the potential inside the electrolyte, which will provide a larger driving force for the dendrite formation. Their proposed solution for suppression of Li dendrites was to lower the electronic conductivity of the SSEs. While limiting the defects and pores in the SSE and lowering the electronic conductivity may suppress the growth and infiltration of Li dendrites, the role of the Li−SSE interface is equally important in dealing with the problem. Wu et al. elucidated the role of a stable interphase layer at the Li-LLZO and Li-LATP interfaces.84 Their findings suggest that formation of an electronically insulating but ionically 22037


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Figure 7. (a) HRTEM image of LCO particles coated with LiI−Li4SnS4 layer through solution process. (b) Nyquist plots and (c) rate-capability response of coated LCO and manually mixed electrodes. Panels a−c reprinted with permission from ref 118. Copyright 2016 Wiley. (d) LZO coating on NCA cathode. (e) EIS spectra for samples with different amount of LZO coating compared to noncoated samples. (f) Cyclic response for coated and bare samples for 100 cycles. Panels d−f reprinted with permission from ref 123. Copyright 2014 Elsevier.

an interface resistance of 13.7 Ω cm2 or below will be required to ensure no Li dendrite formation. Tsai et al.85 found that Li dendrite formation was mostly due to inhomogeneous current distribution at the Li and SSE interface arising from the uneven contact. For improving the contact, they polished the SSE surface and coated with thin Au as buffer layers (Figure 6b). This helped in homogenizing the Li flux that resulted in a significantly lower interface resistance and consequently suppression of Li dendrites. In another work,79 a direct correlation between SSE microstructure and interfacial properties of Li-LLZO cell were demonstrated. It was suggested that SSEs with small-grained surface microstructure perform much better when interfaced with Li. The enhanced performance was attributed to a larger area fraction of low resistivity grain boundaries, which helped in regulating the ionic current densities at the interface. Xu et al. speculated that initiation and propagation of Li dendrites could be suppressed by adding a third phase that is able to react with Li at the interface as well as in the bulk of the SSE.112 They mixed glassy Li3PO4 with LLZO and pressed into pellets. Li3PO4 was able to in situ react with the plated Li at the interfaces to form Li3P during charge cycle (Figure 6c). The authors argued that the reaction is self-limiting and hence imparts interfacial stability. However, they did not comment on the electronic conductivity of Li3P layers that may have an impact on the interfacial stability. An improper contact between Li and SSE due to contaminations and defects at the SSE surface will hinder the ionic transport at the interface. The Li ions will favor flowing in the regions with lower resistance and hence concentrate only at the (limited) points where Li is directly interfaced with the SSE. This will build up an electric field at such sites, which may result in inhomogeneous electrodeposition of Li and ultimately initiate the formation of Li dendrites. Therefore, it is important to realize defect-free and intimate contacts between Li and SSE for suppressing initiation of Li dendrites in SSEs.

5. CATHODE INTERFACE The cathode interface with SSE is technically more challenging than the Li interface. First, because conventional cathodes are composites based on active materials, conductive additives ,and binders, the SSEs have to deal with additional interfaces that will complicate the system further. Second, the cathode materials undergo (de)lithiation, where depending on the type of cathode material, volumetric expansion of varying degrees may occur. This may have a significant impact on the mechanical properties of the interface. For example, in conversion cathodes, the challenge of ensuring intimate contact with SSE during cycling is far bigger than in layered metal-oxide cathodes. Further, because of the crystalline nature of the cathode materials, additional factors come into consideration such as lattice mismatch with SSE and space charge regions formed because of the deficiency of cations/ anions at the interface.113 5.1. Metal-Oxide Cathode Interface with SSE. Cathode Surface Coating. The integration of conventional metal oxide cathodes, especially LCO with sulfide SSE, is challenging, as formation of cobalt sulfide layers takes place at the interface because of mutual diffusion of the electrode−electrolyte species.68,114,115 Cathode surface coating with materials that are stable in a wider range of applied potentials has been demonstrated as an effective strategy to stabilize the cathode− SSE interface. A theoretical study13 reports that Li4Ti5O12, LiTaO3, LiNbO3, Li2SiO3, and Li3PO4 are suitable as coating materials for the LCO cathode because of their chemical stability in a wide potential range (2−4 V). Ohta and coworkers coated LiNbO3 layers on LCO powder by spraying an ethanol solution of alkoxide of Li and Nb with a subsequent heat treatment step.116 The cathode was integrated with a Li3.25Ge0.25P0.75S4 and Li to build a full cell. A much improved performance for the coated samples was observed, attributed to the better electrochemical stability and ionic conductivity of LiNbO3. LiNbO3 coating has also shown to enhance the 22038


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Figure 8. (a) SEM images of the LCO-NASICON SSE annealed at (1) 500, (2) 600, and (3) 700 °C. (b) RBS spectra for the samples annealed at 500, 600, and 700 °C. Panels a, b reprinted with permission from ref 125. Copyright 2017 American Chemical Society. (c) Cross-section SEM image and elemental mapping of the cathode composite coated on LLZO pellet after sintering at 700 °C. Reprinted with permission from ref 126. Copyright 2018 Elsevier Publications. (d) SEM images of LCO-SSE interface after 100 charge/discharge cycles. (e) Nyquist plots after 1st and 100th cycle. (f) 1st and 100th charge/discharge profile for LCO-LLZO-Li cell. Panels d−f reprinted with permission from ref 127. Copyright 2012 Elsevier.

performance of LiMn2O4.117 In another work, Al2O3 layers were formed by the ALD method.68 The layer suppressed the diffusion of Co species from the LCO cathode compared to the uncoated samples that resulted in a much thinner interphase layer. The modification resulted in much better capacity retention for the Al2O3-coated samples, as 90% of the discharge capacity was retained after 25 cycles compared to 70% of discharge capacity for the uncoated samples. Park et al.118 adopted a solution-based approach, where LCO particles were coated with a nanometer thick LiI− Li4SnS4 (Figure 7a). The solution process enabled much better coating of the cathode active materials than the conventional manual mixing approach, which was confirmed by EIS analysis (Figure 7b). The improved contacts had a noticeable effect on the rate performance of the samples where solution coated LCO showed much higher capacity retention than manually coated LCO, especially at higher C rates (Figure 7c). In other work, LCO was coated with LiTaO3, and Li2SiO3 to protect it from reacting with the SSE that resulted in enhanced electrochemical performances.119,120 LiNi1/3Mn1/3Co1/3O2 powders were coated with ZrO2 and LiAlO2 to suppress interfacial resistance when in contact with amorphous Li3PS4.121,122 ZrO2 coating acted as a buffer layer, where it suppressed the undesirable electrode−electrolyte reaction. The samples with ZrO2 coating showed an initial discharge capacity of 115 mAh g−1 and capacity retention for 50 cycles at a current density of 0.1 mA cm−2.121 Ito et al. showed that coating Li-

Ni0.8Co0.15Al0.05O2 (NCA) cathode with Li2O-ZrO2 (LZO) results in improved interface stability between cathode and the sulfide SSE.123 The LZO coated layer was ca. 8 nm in thickness (Figure 7d). A full cell based on graphite anode, sulfide SSE and LZO-NCA was investigated through EIS, where 0.5 mol % LZO samples showed approximately one-quarter less interface resistance than uncoated samples (Figure 7e). The decrease in the interface resistance was attributed to the nanosized LZO layer coating, which protected the active materials from reacting with SSE. Further, the cell retained ∼90% of the initial discharge capacity after 100 cycles as shown in Figure 7f. Cathode−Solid−Electrolyte Annealing. In the case of oxide SSEs, mutual annealing of cathode and solid electrolyte materials is an effective strategy to improve cathode-electrolyte contact. Goodenough and his team prepared a composite of LCO and LLZO materials and studied their electrochemical properties at various annealing temperatures in a quest to improve their physical contact while avoiding chemical reactivity.124 They found out that sintering improves the physical contact between LCO and LLZO. However, they also observed formation of less conducting LLZO tetragonal phase at the interface. Kim and co-workers suggested that to obtain better electrode−electrolyte contacts while avoiding undesirable interface reactions, optimization of the annealing step is crucial.125 They found no reaction between a sputtered LCO thin film and NASICON SSE until 500 °C. However, at higher annealing temperatures (600 and 700 °C), thin interphase layers started to appear at the electrode−SSE junction (Figure 22039


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Figure 9. (a) Charge−discharge voltage profiles for S-CR-LGPS composite cathode for 50 cycles. (b) Schematic representation of S-CR composite cathode and LGPS SSE and the proposed synthesis route. Panels a, b reprinted with permission from ref 135. Copyright 2018 American Chemical Society. (c) Synthesis approach of Li2S−Li6PS5Cl−carbon composite cathode. (d) TEM image and elemental mapping of S and other elements in the nanocomposite. (e) HR-TEM image of the Li2S−Li6PS5Cl−C nanocomposite (inset shows the EDS spectra at point 1 and point 2 marked in the figure). Panels c−e reprinted with permission from ref 138. Copyright 2016 American Chemical Society.

remained intact even after undergoing 100 charge/discharge cycles (Figure 8d). The interface resistance before and after cycling remained almost the same and the cell demonstrated a specific capacity of ∼130 mAh g−1 for 100 cycles (Figure 8e, f). To improve the electrode−electrolyte contact, V2O5 cathode was melted at 800 °C on LLZO surface by a rapid thermal annealing step.128 The cathode−SSE interface resistance decreased to 71 Ω cm2 at room temperature and to 31 Ω cm2 at 100 °C. The full cell maintained a stable charge/ discharge capacity with 97% Coulombic efficiency when tested at 100 °C. The improved performance was attributed to the formation of a continuous and intimate contact between V2O5 cathode and the SSE due to the rapid thermal annealing. 5.2. Sulfur Cathode Interface with SSE. Sulfur (S) is an attractive cathode material for LIB because of its high theoretical capacity of ∼1675 Ah/kg. However, a major issue associated with S cathodes in liquid systems is the formation and dissolution of intermediate polysulfides, which shuttle through the liquid electrolyte, severely degrading the cyclic performance of the cell. SSEs could potentially solve this problem by mechanically blocking the polysulfide shuttle.129 However, maintaining an intimate contact between S and SSE

8a). The layers were amorphous in nature and possibly developed because of interdiffusion of LCO and the SSE species. This is confirmed by Rutherford backscattering spectroscopy (Figure 8b), where no change is observed between samples annealed at 500 °C and the as-deposited samples in terms of Ge and Co concentration at the interface. However, samples annealed at 600 and 700 °C show a drop in Ge and Co species, which suggests interdiffusion of the electrode−electrolyte species at the interface. Han et al. made Li2.3C0.7B0.3O3 react with inherent Li2CO3 layers on LLZO and L C O at h i g h t e m p e r a t u r e s (7 0 0 ° C ) t o f o r m Li2.3−xC0.7+xB0.3−xO3 (LCBO) interphase layers.126 Figure 8c shows a cross-sectional SEM image and elemental distribution in the cathode composite coated on LLZO pellet after the sintering step. LCO and LLZO particles are uniformly distributed in the LCBO matrix, which suggests good reactivity of the materials. The strategy helped in overcoming cathode− SSE interfacial resistance and resulted in a stable electrochemical response. In another work, LCO layers were deposited by pulse laser deposition technique on Li6.75La3Zr1.75Nb0.25O12 and annealed at 600 °C.36 The sputtered layer formed good contact with the SSE, which 22040


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Figure 10. (a) Schematic diagram of the porous/dense Li garnet structures. Porous structures are coated with CNT and S is melted in the pores. (b) Initial charge/discharge profiles and (c) cyclic response for Li−S battery. Panels a−c reprinted with permission from ref 136. Copyright 2017 Royal Society of Chemistry. (d) Schematic diagram and digital photo of Li garnet textile. (e) Digital photo of Li garnet textile sintered on top of dense Li garnet layers; SEM image shows the infiltrated S on the textile pores; elemental distribution in the Li garnet textile; selected charge/ discharge profiles for the Li−S battery based on the textile garnet structures. Panels d, e reprinted with permission from ref 139. Copyright 2018 Elsevier.

In another work, reduced graphene oxide (rGO) was conformally coated with S particles (∼2 nm) to prepare rGO@S, which was further ball-milled with LGPS SSE and conductive carbon to form composite cathodes.129 The cells demonstrated good electrochemical performance, comparable to that of liquid system, when tested at 60 °C. Conformal coating of S on rGO and its uniform distribution in the LGPS matrix provided uniform room to the material’s volume expansion during lithiation, which relieved stress/strain on the S-LGPS interface contact. Similarly introducing S into the pores of mesoporous carbon matrix helped in improving the performance of the battery based on S cathode and thioLISICON SSE.137 Han et al. claimed that the mechanical properties of the interface between lithium sulfide (Li2S) cathode and Li6PS5Cl SSE could be significantly improved if the nanosized active materials and SSE particles can be in situ grown in the carbon matrix.138 They synthesized Li2S− Li6PS5Cl−C composite cathode through a solution-based approach, as shown in schematic Figure 9c. EDS elemental mapping of the composite showed uniform distribution of S, C, and Cl in the structure (Figure 9d). The HR-TEM image in

is a major challenge because of the enormous volume expansion of S (as high as 80%) upon conversion reaction with Li (S to Li2S).130 In the past, various strategies have been proposed to address the problem of poor contact between S and SSEs.131−134 Among them, utilizing carbon-based materials as host matrices129,135 or three-dimensional SSEs architecture for S active materials have been demonstrated as effective strategies.136 Carbon-Based Host Matrix for Cathode Materials. Suzuki et al. used a carbon replica (CR) structure as host for S particles and then adopted a two-step mixing involving initially liquid phase mixing of S/CR in THF solvent, then drying of the chemicals, and finally mechanically ball-milling them.135 The synthesis route is shown in Figure 9b. The cell composed of Li−In alloy, Li10.05Ge1.05P1.95S12 SSE and S-CR-LGPS composite electrode and tested under a constant pressure of 213 MPa, demonstrated discharge capacity of ∼1500 mAh g−1 with ∼100% Coulombic efficiency for 50 cycles (Figure 9a). The high discharge capacity and Coulombic efficiency was attributed to the enhanced contact between S, carbon and SSE, realized through combination of liquid and mechanical mixing. 22041


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Figure 11. (a) Comparison of Li stripping/plating voltage profiles for cells based on LATP SSE with and without LE at the interface. Reprinted with permission from ref 143. Copyright 2018 Elsevier. (b) Schematic diagram highlighting interaction of n-BuLi superbase with LLZO SSE; discharge capacity for the cell for 400 cycles at 100 and 200 μA cm−2. Reprinted with permission from ref 144. Copyright 2011 American Chemical Society. (c) Schematic highlighting the full cell assembly containing LE on either sides of the LATP SSE; Selected charge/discharge voltage profiles (1st to 150th cycles) and discharge capacity for 150 cycles. Reprinted with permission from ref 145. Copyright 2015 Royal Society of Chemistry. (d) Schematic figure outlining cell assembly with LE between sulfur and LLZO; cyclic response of the cell delivering ∼800 mAh g−1 of specific discharge capacity for 500 cycles at a 1 C rate. Reprinted with permission from ref 146. Copyright 2018 Elsevier.

LLZO was synthesized in the form of flexible textile architecture using a template approach (Figure 10d).139 The LLZO textile possessed high porosity and surface/volume ratio. The textile architecture was filled with PEO polymer to form a free-standing composite polymer electrolyte. The room temperature ionic conductivity was on the order of 1 × 10−4 S cm−1 with stable stripping and plating voltage profiles at current densities of 0.2 mA cm−2. To demonstrate a Li−S cell, garnet textiles were sintered onto a dense garnet pellet to form a 3D framework that was later infiltrated with S slurry (S loading ca. 11 mg cm−2) and Li metal was used on the other side of the dense garnet pellet. The cell delivered a high capacity of 1000 mAh g−1 after 40 cycles as shown in Figure 10e. Bruce and his team fabricated NASICON-type Li1.4Al0.4Ge1.6(PO4)3 with bicontinuous 3D micro channels filled with various types of nonconducting polymers through a 3D template approach.140 The SSEs with gyroidal microstructure and epoxy polymer showed ionic conductivity of ∼1 × 10−4 S cm−1 at room temperature. Such 3D designs were proposed to have the potential to attain maximum electrode− SSE contact.

Figure 9e revealed distribution of nanometer-sized Li2S and Li6PS5Cl in the carbon matrix. The solid-state battery composed of Li2S−Li6PS5Cl−C cathode, 80Li2S-20P2S5 SSE, and Li−In alloy anode delivered reversible capacity of 830 mAh g−1 for 60 cycles at a current density of 50 mA g−1. Three-Dimensional Architecture. For fabricating S and oxide SSE composite cathodes, it is important to come up with innovative structural designs capable of overcoming interfacial issues. As mentioned earlier, S undergoes an enormous volume change during (de)insertion. With the brittle nature of oxide SSEs, such volume changes cannot be accommodated, which may lead to mechanical loss of the contacts. An interesting structural concept was presented in a work by Fu and coworkers, where porous LLZO SSE layers prepared by a colloidal technique and precoated with CNT, were loaded with active material (S nanoparticles were melted in the structures) to achieve maximum electrode−electrolyte contact (Figure 10a).136 Later, the composite electrodes were sintered on top of a dense LLZO layer to form a bilayer framework. The Li−S solid-state cell delivered an initial reversible capacity of 645 mAh g−1 at 0.2 mA cm−2 and Coulombic efficiency of ca. 99% (Figures 10b, c). A similar approach was implemented where 22042


LIPON-LCO Li7−3xAlxLa3Zr2O12−Li Li1+x+yAlyTi2−ySixP3−xO12-LCO

Li7La3Zr2O12-LCO Li1.3Al0.3Ti1.7(PO4)3-LCO

interface

SAM

DR-XAS

XPS

LATP-GC-interlayer-LCO (interlayer: NbO2, ZrO2, and MoO2) Li6PS5Cl-LCO, Li6PS5ClLiNi1/3Co1/3Mn1/3O2, Li6PS5ClLiMn2O4

Li10GeP2S12−Li (in situ) LLTO-Li (in situ) Li-β-Li3PS4−Au Li−Li10GeP2S12−Au Li7La3Zr2O12-LCO

Li2S−P2S5−Li (operando)

TEM/EELS coupled with EH spatially resolved Li1+x+yAlyTi2‑ySixP3‑xO12−Li EELS in TEM mode TOF-SIMS Li1.15Y0.15Zr1.85(PO4)3-Li

STEM/EELS

technique analysis revealed an interphase layer with thickness in nm range consisting of intermediate La2CoO4 species no interphase layer observed until temperature ∼500 °C; interlayers started to appear at higher temperatures possibly because of interdiffusion of electrode/electrolyte species disordered LCO resistive layer observed at the electrode−electrolyte interface due to the chemical instability of highly delithiated LCO with LiPON formation of a few nanometer thick tetragonal interphase at the Li-LLZO junction, which helped in preventing further decomposition of the SSE potential variations at the electrode−electrolyte interface due to ionic transport were observed and quantified during cathode redox reaction (Co3+ ↔ Co4+) reactions during charge/discharge cycles, the concentration profiles of Li as well as changes in the crystal and electronic structures of the other elements were simultaneously observed Li1.15Y0.15Zr1.85(PO4)3-Li interface was analyzed to identify the species in the interphase layer responsible for high interface resistance and Li dendrite formation decomposed products such as Li2S and Li3P were observed at the interface under bias conditions; further, because of the presence of oxygen, Li3PO4 phase was formed that subsequently led to Li2O formation interphase layer consisting of decomposed products (Li2S, Li3P, and Li−Ge alloys) was revealed that led to increased interface resistance interphase layer formed between Li and LLTO that was composed of Ti3+, Ti2+ reduced species, and Ti metal continuous breaking/rebuilding of the SSE framework is observed upon Li (de)insertion at the interface at temperatures >300 °C, significant cation interdiffusion along with decomposed products (La2Zr2O7, Li2CO3) were observed at the interface that resulted in higher resistance changes in electronic structures and chemical states of electrode−interlayer−electrolyte interface were studied; NbO2 as interlayer showed less chargetransfer resistance by restricting Co−O bond change and relieving electrode volumetric strain upon (de)-lithiation SAM imaging helped in getting evidence on the spatial separation of LiCl particles and sulfur-containing species (S, Li2Sn, and P2Sx) formed because of oxidation of the SSE with various electrodes

information

Table 2. Summary of Various in Situ and Operando Techniques Used for Characterization of the Interface Region in All-Solid-State Batteries

158

157

154 69 155 156

153

152

151

147 148 149

74 125

ref

ACS Applied Materials & Interfaces Review

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6. INTERFACE WETTING WITH LIQUID ELECTROLYTES Because of the sluggish transport of some of the cathode materials and poor interface kinetics, recent works87,98,141−146 have adopted strategies of adding a small amount of organic liquid electrolyte (LE) at the electrode and inorganic SSE interfaces to improve ionic transport. The LE fills the voids at the interface hence ensures better contact between the SSE and the cathode. Further, LE can penetrate through the pores of the electrode material that helps in accessing more electrode surface area and provides facial paths for ionic transport at the interfaces. Busche et al.142 found a solid−liquid electrolyte interphase (SLEI) layer that contained decomposed products of the liquid and solid electrolytes. They observed that the SLEI layer was of finite thickness and imparted an additional resistance to the system. However, despite the added resistance various studies have shown significant improvement in the interfacial properties due to the addition of LE. Wang et al. discussed the effects of organic LE on the interfacial properties of LiFP-LATP-Li hybrid battery.143 By adding 2 μL of LE (1 M LiPF6 in EC:DMC:DEC, volume ratio 1:1:1) on either side of the SSE, the interfacial resistance decreases to 90 Ω, which is ∼50 times lower than the initial resistance (4470 Ω), which helps obtain stable voltage polarization curves during Li stripping/plating (Figure 11a). Further, a specific capacity of ca. 125 mAh g−1 at a 1 C rate was achieved that decayed by only 8% after 500 cycles. In another work,the role of nbutyllithium (n-BuLi) superbase was emphasized in stabilizing the interface between Li7La3Zr1.5Ta0.5O12 SSE and carbonatebased LE.144 By adding n-BuLi to the conventional LE, much less interface resistance was observed, resulting in improved electrochemical response. The authors claimed that n-BuLi retards the decomposition of LE, suppresses Li+/H+ exchange between LLZO and ambient environment, and offers a favorable medium for ionic transport by lithiating electrode− SSE interface. The schematic of the full cell based on LFP, LE with n-BuLi, LLZO SSE, and Li is shown in Figure 11b. The cells retained 86% percent of the initial capacity operating at 100 μA cm−2, whereas 99% operating at 200 μA cm−2, although capacity decreased significantly (ca. 85 mAh g−1) after switching to 200 μA cm−2. Wing et al. demonstrated a hybrid Li-sulfur battery based on Li2S, LATP and Li with organic LE on either sides of the SSE (Figure 11c)·145 The role of LE was to provide medium for ionic transport at the solid−solid electrode−electrolyte interface. Further, LATP blocked the intermediate polysulfides at the interface hence avoided the shuttle effect. The cell demonstrated a high capacity of ca. 900 mAh g−1 for more than 150 cycles with Coulombic efficiency of ∼ 100% (Figure 11c). In a similar work,146 a highly stable cyclic performance for 500 cycles at a 1 C rate and a high capacity of ca. 800 mAh g−1 of specific discharge capacity was demonstrated by a hybrid battery concept with Li−Au alloy as anode, Li6.4La3Zr1.4Ta0.6O12 as SSE, and P2S5/Li2S as catholyte (Figure 11d). The role of P2S5 was particularly emphasized in enhancing the solubility of Li2S and forming Li3PS4-based interphase layers, which offered improved ionic conductivity. Using other Li-salts, for example, bis(trifluoromethane)sulfonimide (LiTFSI)87,136 and LiCF3SO3,98 in 1,2-dimethoxyethane (DME) and 1,3-dioxolane (DOL) solvents at the sulfur-LLZO interface also demonstrated good electrochemical performances.

7. ELECTRODE−ELECTROLYTE INTERFACE CHARACTERIZATION The interface characterization in ASSBs is very challenging compared to their liquid counterparts because of the physical and electrochemical complexities associated with the solid− solid interface. In situ and/or operando analysis is preferred to provide true and in-depth information on the formation and evolution of the interphase layers during battery cycling. In recent times, the dynamic conditions at the electrode−SSE interface have been intensively explored using advanced characterization techniques.147−158 Table 2 summarizes some of those techniques where interfaces between Li, cathodes, and SSEs were studied. In situ scanning transmission electron microscopy (STEM) and electron energy loss spectroscopy (EELS) showed formation of interphase layers in LLZO-LCO, LATP-LCO, and LIPON-LCO interface.74,125,147,148 By coupling TEM and EELS characterization with in situ electron holograph (EH), electric potential distributions at the Li and LCO interfaces with Li1+x+yAlyTi2−ySixP3−xO12 SSE were obtained.151 Also, by adopting spatial resolved spectroscopy variations in the Li concentration profiles and changes in the crystal and electronic structures of the other elements of Li1+x+yAlyTi2−ySixP3−xO12 were simultaneously observed during charge/discharge cycles.152 Time-of-flight secondary-ion mass spectrometry (TOF-SIMS) was employed to qualitatively analyze the 3D composition of the Li1.15Y0.15Zr1.85(PO4)3-Li interphase layer to uncover the factors responsible for high interface resistance and Li dendrite formation.152 X-ray photoelectron spectroscopy (XPS) studies were also conducted to get quantitative and qualitative information on the reaction products when Li is interfaced with LLTO, Li 2 S−P 2 S 5 , LGPS, and βLi3PS4.69,153−155 Depth-resolved X-ray absorption spectroscopy (DR-XAS) was employed to get insights into the LATPLCO interface at nanoscale. The study revealed that presence of an interlayer (NbO2) suppresses the interface resistance by restricting Co−O bond change and relieving electrode volumetric strain upon delithiation.156 Scanning auger microscopy (SAM) was conducted on interfaces between various conventional cathodes (LCO, LiNi1/3Co1/3Mn1/3O2, LiMn2O4) and Li6PS5Cl.158 The SAM analysis was particularly helpful in getting information on species formed due to oxidation of the SSE, which were not detected by conventional microscopy. For example, spatial separation of LiCl and S containing species (S, Li2Sn, and P2Sx) suggested formation of an interphase layer. The aforementioned in situ and operando techniques assist in better understanding of the structural, compositional, and morphological variations occurring at the electrode−SSE interfaces, which may be correlated with the capacity degradation of the cells. Further, improved diagnosis will allow researchers to devise better strategies in dealing with the underlying problems related to the electrode-SSEs interfaces. 8. CONCLUSION AND OUTLOOK All-solid-state battery based on inorganic SSE is a potential candidate for EV application because of its widely acknowledged safety benefits, energy density, and thermal stability. However, their employment in practical systems is challenging. Although there has been significant progress in achieving SSE’s ionic conductivities, which are now on a par with liquid electrolytes, the issue of high interface resistance at the solid− 22044


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operando characterization techniques, which could provide insights into the formation of interphase layers and the transport phenomenon happening at the interfaces.

solid electrode−SSE interface limits their practical application. Various factors responsible for the interfacial resistance include chemical incompatibility, electrochemical instability, and mechanical issues between the materials. In this review, we focused on the Li and cathode interfaces in oxide and sulfide SSEs and summarized recent reports dealing with the interfacial processes that take place at the electrode−SSE interfaces in such systems, how they impact the performance of the ASSB, and what the strategies are to address these issues. At the electrode−SSE junction, interphase layers should be electronically insulating so that passivation of the SSE is achieved which may help in bridging the potentials of the electrodes and the SSEs hence providing kinetic stability to the charge transport. Depending on the type of SSE (oxide or sulfide), the electrode−electrolyte interfaces offer different challenges. Although oxide SSEs have generally good electrochemical stability with Li and other cathode materials, their brittle and stiff nature presents mechanical issues at the interface. Also, poor Li wetting toward oxide SSEs due to presence of surface carbonate and hydroxide layers is not only responsible for large voltage polarization but also initiates the formation of Li dendrites at the Li−SSE interfaces. Therefore, efforts are directed toward improving the electrode-SSE contact by various strategies such as enhancing electrode reactivity with the SSE at the interface with nanolayer coatings, surface engineering of the SSE, and sputtering or melting of the electrode materials on the SSEs. Further, wetting the interfaces with LE or using thin flexible polymeric interlayers at the interface helps in reducing interface resistance. Sulfide SSEs, on the other hand, suffer from chemical instability with Li metal (sulfide SSEs decompose against Li) and metal-oxide cathodes (formation of mutual-diffused layers at the interface), which results in a thick interphase layer that degrades the ionic transport at the interface. Strategies that are adopted to achieve interface stability include use of Li−In alloys as anodes and surface coatings of cathode with materials that are electrochemically stable with sulfide SSEs. Further, carbon-based materials serving as host are effective in relieving mechanical stresses during (de)lithation of the electrode active materials especially those undergoing enormous volume change. Along with innovative strategies to address the issues of interfaces in solid-systems, utilization of various advanced characterization techniques such as STEM/EELS, TEM, XPS, TOF-SIMS, DRXAS, and SAM are equally important for deep understanding of the complex and dynamic electrochemical processes happening at the electrode-SSE interface. In the future, dedicated research efforts are required to address the issue of high interfacial resistance in ASSBs. To start with, the SSE materials should have favorable physical and (electro)chemical properties. For example, thermodynamically the SSEs should be stable with Li and other promising cathodes in a wide potential range. Mechanically, they should have a deformable/ductile nature, which will simplify their device integration. It is important to understand the mechanism of the ionic transport at the electrode−SSE interfaces. Once the mechanism is fully understood and the issues are properly identified, a wide range of solutions including interface modification via coating layers of stable materials, polymeric/gel interlayers, and 3D designing/grading can be suggested. In this regards, computational studies as supplement to experimental efforts may allow researchers to devise better strategies. Lastly, further advancements are needed for development of more powerful in situ and

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AUTHOR INFORMATION

Corresponding Author

*Email: [email protected]. ORCID

Syed Atif Pervez: 0000-0002-7134-7103 Musa Ali Cambaz: 0000-0002-4249-3486 Venkataraman Thangadurai: 0000-0001-6256-6307 Author Contributions

The manuscript was written through contributions of all authors. All authors have given approval to the final version of the manuscript. Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS We gratefully acknowledge financial support by the Alexander von Humboldt Foundation, Bonn, Germany. We also thank the Deutsche Forschungsgemeinschaft for financial support under project ID 422053626 (Cluster of Excellence “Post-Li Storage”). This work contributes to the research in CELEST (Center for Electrochemical Energy Storage Ulm-Karlsruhe). V.T. thanks the University of Calgary, Calgary, Canada, as well as the Natural Sciences and Engineering Research Council of Canada (NSERC), for their support.

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Interface in Solid-State Lithium Battery: Challenges, Progress, and

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