Interface in Solid-State Li Battery: Challenges, Progress and Outlook

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Interface in Solid-State Li Battery: Challenges, Progress and Outlook Syed Atif Pervez, Musa Ali Cambaz, Venkataraman Thangadurai, and Maximilian Fichtner ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.9b02675 • Publication Date (Web): 30 May 2019 Downloaded from http://pubs.acs.org on May 30, 2019

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ACS Applied Materials & Interfaces

Interface in Solid-State Li Battery: Challenges, Progress and Outlook Syed Atif Perveza,*, Musa Ali Cambaza, Venkataraman Thangaduraib, Maximilian Fichtnera a

Helmholtz Institute Ulm, Helmholtzstraße, 11, 89081, Ulm, Germany

* Corresponding Author; E-mail: [email protected] b

University of Calgary, Department of Chemistry, 2500 University Dr NW, Calgary, AB,

T2N 1N4, Canada; E-mail: [email protected]

KEYWORDS: interface resistance, solid-electrolyte, Li dendrites, solid-state battery, interfacial characterization

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 chemo-mechanical 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.

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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 In order to achieve this goal, development of technologies with ideally no or reduced greenhouse gas emission is needed for grid-electrification and automotive transportation. Renewable energy storage using Li-ion battery (LIB) is a viable option to achieve the aforementioned goal. In the last 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 solventbased 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 high-energy applications such as EVs and portable electronics is limited due to 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 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

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addressed by replacing liquid electrolytes with solid-state electrolytes (SSEs). SSEs are considered safer due to their non-flammable 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.12,13,14 Oxides include Li-garnets,15,16 perovskites,17,18 Li Super Ionic Conductors (LISICONs),19,20,21 and Na Super Ionic Conductors (NASICONs).2,22

Sulfides

include

thio-LISICON-type

LGPS23,24

and

glassy-type

Li2S−P2S5.25,26,27 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 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. While 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). 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

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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 reviewed recent works, which investigate the role of high Li-SSE interface resistance in initiating Li dendrites, and on hybrid liquid-solid systems for advanced batteries. We also reviewed in situ and operando characterization techniques currently used to study the dynamic processes happening at the electrode-SSE interfaces.

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 towards solving them are listed.


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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 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 plat 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: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.,



E A  exp( a ) T kT

(1)

where ‘’ is the conductivity of the electrolyte in S cm-1, ‘A’ the frequency factor, ‘Ea’ activation energy expressed in joules per mole, ‘k’ is Boltzmann’s constant and ‘T’ temperature measured in Kelvin. In order to build a working battery the integration of the SSE with various electrodes should be chemically stable at wide range of potentials.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 de-merits. 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 room temperature ionic conductivity of 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 (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. Li rich 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+, 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 perovskite-type 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 10−3 S cm-1 at room temperature with activation energy of ~ 0.3eV. Highly conductive Ti-based perovskite SSEs are chemically stable in air/humid conditions and in a 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


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negative electrodes operating at low voltages. 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 group. Early reports from P.G. Bruce and A.D. 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 (~ 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 oC.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 one-dimensional along the c-axis and twodimensional in ab plane with 0.16 and 0.26 eV activation energies, respectively.42 These sulfides offer high ionic conductivity in comparison to oxides due to low grain boundary resistance. Also, they are mechanically robust due to their ductile nature. However, they are costly due to presence of Ge, therefore other metals such as Sn and Si are employed in order 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) which increase the interface resistance.13 Also,



they are highly moisture sensitive and possibly hazardous due to formation of H2S gas if exposed to air for long periods of time.44 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 ~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 70-80) is of particular interest because of high ionic conductivity (10-3-10-4 S cm-1).50,51 Such sulfide glass-ceramic 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 F4̅3m space group) where P atoms co-ordinate with S to


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form PS4 tetrahedrons while Li ions are distributed over the tetrahedral interstices (48h and 24g sites).56 DFT calculations showed that higher 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

environment


is

a

major

concern.

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Table 1. Summary of the crystal structures, ionic conductivity, activation energy, performance and limitations of various inorganic solidelectrolytes

Class

Material/Crystal structure

σLi+(S cm-1) at 25 oC

Ea(eV)

Performance

Oxides

Garnet-type Li7La3Zr2O12

10-3–10-4

0.35–0.4

Thermodynamic stability vs. Li at wide potential range; higher mechanical strength; thermally stable at higher temperatures

Perovskite (Li3xLa2/3xTiO3)(x=0.1)

~10-3

0.3–0.4

Stable in air and a wide temperature range (4K-1600K)


Limitations

Ref.

Brittle nature; formation of insulating carbonates and hydroxides in humid environments

15,16, 19

Reduces vs. Li at potential ~ 1.5 V

14,17, 18

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Sulfides


LISICON-type Li2+2xZn1−xGeO4

10-7–10-3

0.4–0.6

Thermally stable at high temperatures

Low ionic conductivity; thermodynamically not stable with Li

19,20, 21

NASICON-type LiM2(XO4)3(M= Ge,Ti, Zr; X= S,P,As )

~10-3

0.3–0.4

High bulk ionic conductivity

Reduces (Ti4+ to Ti3+) when in contact with Li metal

2,22

Li10MP2S12 (M= Si, Ge, Sn)

~10-2

0.2–0.25

High total ionic conductivity; mechanically robust due to deformable nature

Relatively low potential window; costly due to Ge metal; high moisture sensitive; hazardous due to formation of H2S gas when exposed to water

23,24


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Argyrodites

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Li2S−P2S5 (Li7P3S11)

10-2–10-3

0.38–0.40

High ionic conductivity; improved stability with Li than LGPS; ductile in nature

Sensitive to humid environments

25,26, 27

Li6PS5X (X= Cl, Br, I)

10-2–10-3

~ 0.3

Very high ionic conductivity

Not stable in air and humid environment; decomposes vs. Li metal

28,54, 58,59


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3. Interfacial phenomena at electrode solid-electrolyte interface Understanding the mechanism behind the interfacial processes in solid-state batteries is challenging due to the complex nature of the interface between electrode active materials and the SSEs. The origin of interfacial phenomenon in solid-system is determined by two important factors, firstly, the chemical and electrochemical compatibility between the electrodes and SSEs and secondly, 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 due to the continuous change of 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. ɸLi is the electrostatic potential of Li anode. ɸLixMyO2 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   Li   Lix M y O2

(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 due to 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, 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.65,66,67,68 Ge-based sulfide SSEs such as LGPS are unstable with Li since they start reducing at ~ 1.7 V and oxidizing at ~ 2.15 V leaving 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.


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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. ɸLi is the electrostatic potential of Li and ɸLixMyO2 is the potential of LixMyO2.2 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 than 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 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 V 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 (~ 0.05V 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, Nb-doped 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 due to the formation of a La2CoO4 based interphase layer at the LCO-LLZO junction. Further,in case of LiNi0.5 Mn1.5O4 (LNMO) cathode materials and LLZO, Hansel et al. observed a potential drop at ~ 3.8 V during 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. In terms of mechanical properties, sulfide SSEs have 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 studies 62,76,77 have shown that loss of contact still occur between active materials and the sulfide SSE during cycling. For example, in a LiNi0.8Co0.1Mn0.1O2 solid-state battery


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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 electrochemo-mechanical 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 micro-cracking. 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 sulfidetype.10 4. Lithium anode-SSE interface 4.1. Improving Li-SSE contact with thin coatings The poor physical contact due to 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 them on top of the LLZO pellets have helped in improving the electrode-electrolyte contact.79,80 A better strategy is coating LLZO surface with a thin layer (thickness: 10-20 nm) of materials having good chemical reactivity with Li. Due to 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 and b, 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 form Li-Zn alloys. During alloying process, Li diffuse through the metal-oxide layer hence improves the Li-LLZO contact. Recently, Zhou et



al. coated ZnO on LLZO by 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 Schematic Figure 3c shows the process of lithiation of the ZnO layer, also shown in photos (Figure 3d and 3e). 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). 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 Li-Mg-LLZO interface, the Mg diffuses in to the bulk Li forming Li ― Mg alloys that improved the Li-wettability (Figure 3g and h). The Li-symmetrical cells showed much lower interface resistance and stable Li stripping/plating voltage profiles for 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 Ge88 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 over-potential and ensured low interface resistance. In another work, 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 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 and 4b).92 The work was aimed at innovative architectural designing of the surface with


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micrometer-scale 3D pattern, which could increase areal contact between Li and SSE. To demonstrate

feasibility

of

the

structures,

Li-symmetrical

cells

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. Reprinted with permission.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. Reprinted with permission.83 Copyright 2016, Wiley.



were fabricated (Figure 4b) and cycled reversibly for 150 hours at current density of 0.33 mA cm-2. A similar concept is presented in another work, where a porous-dense-porous 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 over-potential. 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 Li2CO3 and LiOH that can grow on the surface upon exposure to air and humidity.94,95,96 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 oC.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


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modify the surface of the SSEs are considered simple and cost effective in contrast to vacuum based coating techniques.

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. Reprinted with permission.92 Copyright 2018, Wiley (c) SEM images showing porousdense-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 non-polished LLZO garnet pellets. Reprinted with permission.94 Copyright 2017, American Chemical Society. 4.3. Polymeric interlayer at the interface Owing to their flexible and cohesive nature, Li ion conducting polymer layers at the electrodeSSE interface are a good choice to avoid the solid-solid contact between electrode and SSE.



Polyethylene 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 Lielectrolyte 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 NASICON-type SSE by PEO based solid polymer electrolyte (SPE).102 The cell comprising of Li, SPE, SSE pellet and LFP cathode was assembled in a multi-layer form (Figure 5a and b). The PEO interlayer significantly reduced the interface resistance, which resulted in stable Li stripping/plating voltage profiles with much lower voltage polarization. Further, high reversible capacity of 150 mAh g-1, stable for more than 200 cycles at 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 (~ 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 doesn’t remain in the film form rather transforms into a gel which could diffuse across the interphase layer, leading to loss of Li+ from cathode.103 In addition, due to the low transference number (0.2-0.5) of PEO polymers rapid anion depletion can occur at the interface building a large electric field, which may favor decomposition of the electrolyte and enhanced Li dendrite nucleation. To address the issues, they used a cross-linked poly(ethylene glycol) methyl ether acrylate (CPMEA) polymer, thermally stable at much higher temperatures (270 °C) and ionically conductive (~ 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


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the anodic and cathodic interfacial resistance with the SSE to 248 Ω cm2 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).

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. Reprinted with permission.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. Reprinted with permission.104 Copyright 2017, American Chemical Society.

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 owing to their higher shear modulus than polymer SSEs, should be more efficient in suppressing Li dendrite formations. However, various studies79,84,85,106,107,108,109,110,111 have shown formation of Li dendrites in crystalline SSEs which eventually lead to short circuit despite their relatively



high shear modulus than polymer and liquid electrolytes. Porz et al.110 conducted a study on amorphous Li2S-P2S5, polycrystalline β-Li3PS4, 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, irrespective 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 Based on their experimental results, they proposed mechanism for the growth of Li dendrites, which was based on two critical conditions. Firstly, the presence of mobile electrons inside the electrolyte and secondly, 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 suppressing the growth and infiltration of Li dendrites, the role of 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 conducting interphase layer will help in alleviating Li dendrite growth at the interface. They justified their hypothesis by adding 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 LLZO surface by filling the


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microstructures with Si nanoparticles by a facial polishing process. Li on 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 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 due to 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, 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 to flow 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 Li interface. Firstly, since 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. Secondly, the cathode materials undergo (de-)lithiation where, depending on the type of cathode material, volumetric expansion of varying degree may occur. This may have a significant impact on the mechanical properties of the interface. For example, in conversion cathodes the challenge to ensure intimate contact with SSE during cycling is far bigger than in layered metal-oxide


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cathodes. Further owing to the crystalline nature of the cathode materials additional factors come into consideration such as lattice mismatch with SSE and space charge regions formed due to deficiency of cations/anions at the interface.113

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.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.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.112 Copyright 2017, Elsevier.

5.1. Metal-oxide cathodes interface with SSE Cathode surface coating: The integration of conventional metal oxide cathodes especially LCO with sulfide SSE is challenging, since formation of cobalt sulfide layers takes place at the 27 Environment ACS Paragon Plus


interface due to 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 cathode-SSE interface. A theoretical study13 reports that Li4Ti5O12, LiTaO3, LiNbO3, Li2SiO3 and Li3PO4 are suitable as coating materials for LCO cathode due to their chemical stability in a wide potential range (2 ― 4V). 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 performance of LiMn2O4.117 In another work, Al2O3 layers were formed by 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


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ZrO2 and LiAlO2 to suppress interfacial resistance when in contact with amorphous Li3PS4.121,122

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. Reprinted with permission.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. Reprinted with permission.123 Copyright 2014, Elsevier. 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 LiNi0.8Co0.15Al0.05O2 (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 nano-sized 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 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 in order 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 oC. However, at higher annealing temperatures (600 oC, 700 oC), thin interphase layers started to appear at the electrode-SSE junction (Figure 8a). The layers were amorphous in nature and developed possibly due to inter-diffusion of LCO and the SSE species. This is confirmed by Rutherford backscattering spectroscopy (Figure 8b) where no change is observed between samples annealed at 500oC and the as-deposited samples in terms of Ge and Co concentration at the interface. However, samples annealed at 600 oC and 700 oC show drop in Ge and Co species which suggest inter-diffusion of the electrode-electrolyte species at the interface. Han et al. made Li2.3C0.7B0.3O3 react with inherent Li2CO3 layers on LLZO and LCO at high temperatures (700 oC) to form Li2.3−xC0.7+xB0.3−xO3 (LCBO) interphase layers.126 Figure 8c shows cross-sectional SEM image and elemental distribution in the cathode composite


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coated on LLZO pellet after the sintering step. LCO and LLZO particles are uniformly distributed in the LCBO matrix, which suggest 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 oC.127 The sputtered layer formed good contact with the SSE, which 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 and f). In order to improve the electrode-electrolyte contact, V2O5 cathode was melted at 800 oC 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.



Figure 8. (a) SEM images of the LCO-NASICON SSE annealed at (1) 500 oC (2) 600 oC and (3) 700 oC. (b) RBS spectra for the samples annealed at 500 oC, 600 oC and 700 oC. Reprinted with permission.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 oC. Reprinted with permission.126 Copyright 2018, Elsevier Publications (d) SEM Images of LCOSSE 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. Reprinted with permission.127 Copyright 2012, Elsevier. 5.2. Sulfur cathode interface with SSE Sulfur (S) is an attractive cathode material for LIB owing to 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


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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 is a major challenge due to 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.131,132,133,134 Among them utilizing carbon based materials as host matrices

129,135

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 mechanical 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. In another work, reduced Graphene Oxide (rGO) was conformally coated with S particles (~ 2nm) 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 thio-LISICON 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 nano-sized 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 Figure 9e revealed distribution of nano-metric sized Li2S and Li6PS5Cl in the carbon matrix. The solidstate 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.


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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. Reprinted with permission.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 nano-composite. (e) The HR-TEM image of the Li2S−Li6PS5Cl−C nano-composite (inset shows the EDS spectra at point 1 and point 2 marked in the figure). Reprinted with permission.138 Copyright 2016 American Chemical Society. 35 Environment ACS Paragon Plus


Three-dimensional architectures: 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 enormous volume change during (de-)insertion and with 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 co-workers where porous LLZO SSE layers prepared by colloidal technique and pre-coated with CNTs, 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 bi-layer framework. The Li-S solid-state cell delivered initial reversible capacity of 645 mAh g-1 at 0.2 mA cm-2 and Coulombic efficiency of ca. 99 % (Figures 10b and c). A similar approach was implemented where 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 in the order of 10-4 S cm-1 with stable stripping and plating voltage profiles at current densities of 0.2 mA cm-2. In order to demonstrate a Li-S cell, garnet textiles were sintered onto a dense garnet pellet to form a 3D framework which 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 non-conducting polymers through 3D template approach.140 The SSEs with gyroidal microstructure and epoxy polymer showed ionic


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conductivity of ~ 10-4 S cm-1 at room temperature. Such 3D designs were proposed to have the potential to attain maximum electrode-SSE contact.

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. Reprinted with permission.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. Reprinted with permission.139 Copyright 2018, Elsevier.



6. Interface wetting with liquid electrolytes Owing to the sluggish transport of some of the cathode materials and poor interface kinetics, recent works87,98,141,142,143,144,145,146 have adopted strategies of adding 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 2µL of LE (1M LiPF6 in EC: DMC: DEC, volume ratio 1:1:1) on either side of the SSE, the interfacial resistance decreases to 90 Ω which was ~ 50 times lower than the initial resistance (4470 Ω) which helped in obtaining stable voltage polarization curves during Li stripping/plating (Figure 11a). Further a specific capacity of ca. 125 mAh g−1 at 1 C-rate was achieved which decayed by only 8% after 500 cycles. In another work, role of n-Butyllithium (n-BuLi) super-base was emphasized in stabilizing the interface between Li7La3Zr1.5Ta0.5O12 SSE and carbonate-based 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


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based on LFP, LE with n-BuLi, LLZO SSE and Li is shown in Figure 11b. 86 % of the initial capacity was retained for cells operating at 100 µA cm-2 while 99% operating at 200 µA cm-2 although capacity decreased significantly (ca. 85 mAh g-1) after switching to 200 µA cm-2.

Figure 11. (a) Comparison of Li striping/plating voltage profiles for cells based on LATP SSE with and without LE at the interface. Reprinted with permission.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 µA cm-2 and 200 µA cm-2. Reprinted with permission.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-150th cycles) and discharge capacity for 150 cycles. Reprinted with permission.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 1 C-rate. Reprinted with permission.146 Copyright 2018, Elsevier.

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).145The 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 1Crate 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 LiCF3SO398 in 1,2-dimethoxyethane (DME) and 1,3-dioxolane (DOL) solvents at the sulfurLLZO interface also demonstrated good electrochemical performances.

7. Electrode-electrolyte interface characterization The interface characterization in ASSBs is very challenging compared to their liquid counterparts due to 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,


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the dynamic conditions at the electrode-SSE interface have been intensively explored using advanced characterization techniques.147,148,149,150,151,152,153,154,155,156,157,158 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, Li2S-P2S5, LGPS and β-Li3PS4.69,153,154,155 Depth-resolved X-ray absorption spectroscopy (DR-XAS) was employed to get insights into the LATP-LCO 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 de-lithiation.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



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microscopy. For example, spatial separation of LiCl and S containing species (S, Li2Sn, and P2Sx) suggested formation of an interphase layer. Table 2. Summary of various in situ and operando techniques used for the characterization of interface region in all-solid-state batteries Technique

Interface

Information

Reference

STEM/EELS

Li7La3Zr2O12-LCO

Analysis revealed an interphase layer with thickness

74

in nm range consisting of intermediate La2CoO4 species Li1.3Al0.3Ti1.7(PO4)3-LCO

No interphase layer observed till temperature ~ 500 oC.

Interlayers

started

to

appear

at

125

higher

temperatures possibly due to inter-diffusion of electrode/electrolyte species LIPON-LCO

A disordered LCO resistive layer observed at the

147

electrode-electrolyte interface due to the chemical instability of highly delithiated LCO with LiPON

Li7-3xAlxLa3Zr2O12-Li

Formation of a few nanometer thick tetragonal

148

interphase at the Li-LLZO junction, which helped in preventing further decomposition of the SSE

TEM/EELS

Li1+x+yAlyTi2-ySixP3-xO12-

Potential variations at the electrode-electrolyte

coupled with EH

LCO

interface due to ionic transport were observed and

149

quantified during cathode redox reaction (Co3+ ↔ Co4+) reactions Spatially-

Li1+x+yAlyTi2-ySixP3-xO12-Li

During charge/discharge cycles, the concentration

resolved EELS in

profiles of Li as well as changes in the crystal and

TEM mode

electronic structures of the other elements were

151

simultaneously observed TOF-SIMS

Li1.15Y0.15Zr1.85(PO4)3-Li

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


152

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dendrite formation XPS

Li2S-P2S5-Li

Decomposed products such as Li2S and Li3P were

(operando)

observed at the interface under bias conditions.

153

Further, due to presence of oxygen Li3PO4 phase was formed that subsequently lead to Li2O formation Li10GeP2S12-Li

Interphase layer consisting of decomposed products

(in situ)

(Li2S, Li3P and Li-Ge alloys) was revealed that lead to increased interface resistance

LLTO-Li

An interphase layer formed between Li and LLTO

(in situ)

which was composed of Ti3+, Ti2+ reduced species

154 69

and Ti metal Li-β-Li3PS4-Au

A continuous breaking/rebuilding of the SSE

Li-Li10GeP2S12-Au

framework is observed upon Li (de-)insertion at the interface

Li7La3Zr2O12-LCO

At temperatures > 300°C significant cation interdiffusion

along

with

decomposed

155 156

products

(La2Zr2O7, Li2CO3) were observed at the interface which resulted in higher resistance DR-XAS

LATP-GC-interlayer-LCO

Changes in electronic structures and chemical states of

(Interlayer:NbO2, ZrO2 and

electrode-interlayer-electrolyte interface are studied.

MoO2)

NbO2 as interlayer showed less charge transfer

157

resistance by restricting Co-O bond change and relieving electrode volumetric strain upon (de)-lithiation SAM

Li6PS5Cl-LCO,

Li6PS5Cl-

SAM imaging helped in getting evidence on the

LiNi1/3Co1/3Mn1/3O2,

spatial separation of LiCl particles and sulfur

Li6PS5Cl-LiMn2O4

containing species (S, Li2Sn, and P2Sx) formed due

158

to oxidation of the SSE with various electrodes

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 due to their 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-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 are the strategies 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. While 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 towards 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


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Li–SSE interfaces. Therefore, efforts are directed towards improving the electrode-SSE contact by various strategies such as enhancing electrode reactivity with the SSE at the interface with nano-layer 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, suffers from chemical instability with Li metal (sulfide SSEs decompose against Li) and metaloxide cathodes (formation of mutual-diffused layers at the interface) which results in thick interphase layer that degrade 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, DR-XAS, and SAM are equally important for deep understanding of the complex and dynamic electrochemical processes happening at the electrode-SSE interface. In 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 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, wide range 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 operando characterization techniques, which could provide insights into the formation of interphase layers and the transport phenomenon happening at the interfaces. Corresponding Author *Email: [email protected] Author Contributions The manuscript was written through contributions of all authors. All authors have given approval to the final version of the manuscript. Acknowledgements 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). One of us, 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 Li Battery: Challenges, Progress and Outlook

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