Designing Ionic Conductors: The Interplay between StructuralPhenomena and Interfaces in Thiophosphate-Based Solid-StateBatteriesSean P. Culver,† Raimund Koerver,† Thorben Krauskopf,† and Wolfgang G. Zeier*,†

†Institute of Physical Chemistry, Justus-Liebig-University Giessen, Heinrich-Buff-Ring 17, D-35392 Giessen, Germany

ABSTRACT: Elucidating the underlying structural principles that governionic transport in thiophosphate solid electrolytes will enable the discoveryof novel ionic conductors. Additionally, improving the properties of ionicconductors and exacting control over interfacial reactions and interphasestabilities are critical to the advancement of solid-state batteries. In thisperspective, we focus on two major aspects at the foundation of solid-statebattery development. First, we address the typical static structural require-ments for achieving high ionic conductivities within thiophosphates, which isthen extended to how a dynamic lattice and local structural effects caninfluence ionic transport. Furthermore, we provide an overview of some ofthe challenges that are currently hindering the progress of solid-state batteryresearch, with particular attention being paid to interfacial instabilities andmechanochemical effects. We hope that this perspective provides a uniqueoutlook on ionic conduction in thiophosphates toward the design of future solid electrolytes and highlights the importance ofinterfacial chemistry in the optimization of solid-state battery devices.

1. INTRODUCTION

Lithium-ion batteries are a ubiquitous energy storage technol-ogy that has transformed the role of commercial electronics insociety. The high energy and power densities, along with thegood reliability and cyclability, have made lithium-ion batteriesan obvious choice for powering our commercial electronicdevices.1 While today’s state-of-the-art batteries offer goodperformances, further improvements to the energy density arenot expected without the use of a lithium metal anode andhigh-voltage cathodes.2 Unfortunately, the commonly usedliquid electrolytes are easily oxidized at higher voltages and arenot able to suppress dendrite formation when lithium metalanodes are employed, leading to a variety of safety concerns.3,4

However, the use of a solid electrolyte (SE) separator maycircumvent the aforementioned problems. Recent efforts in SEdevelopment have provided the field with a variety of excep-tionally fast ionic conductors that can be practically employed.6,7

The use of SEs is thought to provide some notable advantages5

over the current battery technology: (1) the solid separatorshould mitigate unwanted electrode cross-talk of polysulfides ortransition-metal cations,8 while also eliminating typical leakageissues associated with liquid-based architectures. (2) A negli-gible partial electronic conductivity should prevent self-discharge.Moreover, (3) the ionic transference number of nearly unitymeans only lithium ions are moving, so there should not be anypolarization resistance at high current densities.7 (4) Chemicalstability at elevated temperatures should also be improved,while additionally preventing the “freezing out” of the electro-lyte at low temperatures. Finally, (5) the superior mechanicalrigidity may even prevent dendrite formation caused during the

electrodeposition of lithium.9,10 While dendrite formation stillremains a challenge at high current densities and a variety ofelectrolyte-based interfacial instabilities have been reportedagainst metal anodes,3,4,11−15 the use of a lithium metal anodewill be imperative in achieving high energy densities in SSBs.One possible approach for transitioning from a typical lithium-

ion battery architecture to an SSB utilizing a lithium metal anodeis illustrated in Figure 1. While the decreased cell volume andthin metal anode provide gains in gravimetric and volumetricenergy densities, the schematic already alludes to potentialdesign issues. For example, the separator needs to be exception-ally thin, while at the same time preventing short-circuits. More-over, ideal percolation must be achieved to electrochemicallyaccess all of the active material,16 and a high ionic conductivity inthe electrolyte is required to reduce overpotentials when increasingelectrode thicknesses.17 Persistent grain contacts may also leadto interfacial reactions, high charge transfer resistances, andmicrostructural strain.18−20 Ultimately, all of the aforementionedconcerns have led to the predominant usage of high-performancethiophosphates21−23 in SSB research.Nevertheless, despite strong contributions from thiophosphate-

based ionic conductors,6,7,24 the list of materials exhibiting suffi-cient ionic conductivity for practical SSB applications remainsexceedingly short. While a multitude of approaches for enhancingionic conductivity have been developed over the years (e.g.,tuning of the crystal structure,25 elemental substitutions,26,27

Received: March 28, 2018Revised: May 10, 2018Published: May 11, 2018

Perspective

pubs.acs.org/cmCite This: Chem. Mater. 2018, 30, 4179−4192

© 2018 American Chemical Society 4179 DOI: 10.1021/acs.chemmater.8b01293Chem. Mater. 2018, 30, 4179−4192

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and microstructural modifications28), the success of any givenapproach is strongly dependent on our understanding of ionicconduction in solids.29 Thus far, much is already known regard-ing static structural influences on transport behavior; however,recent studies have pointed toward dynamic and local structuraleffects as well.30−32 Therefore, by expanding our knowledge ofthe complex processes governing ionic diffusion in thiophos-phates, we can further enhance relevant transport propertiesand develop more general strategies for effectively tuning ionicmobility in SEs.In this perspective, we provide an overview of contemporary

approaches for understanding and designing lithium thiophos-phate solid electrolytes. First, we focus on the typical structuralrequirements for good ionic conduction while showing thatstructural changes must be monitored to elucidate the conduc-tion mechanisms and effectively tailor the performance ofthiophosphates. We then present some of the recent progresson determining the influence of lattice dynamics on ionic trans-port, and further, we discuss the implications associated withdisagreeing local and average structures within ionic conduc-tors. Finally, we address the current challenges in SSBs con-cerning interfacial instabilities, redox activity of the electrolytes,and mechanochemical changes that occur during cycling.We hope that by better defining the structure−property relation-ships at play in solid ionic conductors, in addition to clarifying the

interfacial reactions and performance limitations that are cur-rently impeding battery development, we can guide futureefforts in the fields of ionic conductors and solid-state batteries.

2. STRUCTURAL INFLUENCES ON IONICCONDUCTION

Solid ionic conductors have been studied for decades, and assuch, have become indispensable for a myriad of applica-tions.1,5,32−34 In particular, these ionic conductors have beenheavily investigated as solid electrolytes (i.e., separators) insolid-state batteries. However, to be truly applicable, the tar-get materials must possess exceptional transport properties.To accomplish this goal, improvements on the existing designprinciples of SEs, from a structural standpoint, need to bemade. In considering how to approach this complicated task,one must develop a strong understanding over the structural influ-ences governing ionic conduction in solids, thereby defining theappropriate routes to success.

2.1. Static Lattice Effects. In general, the ionic conductiv-ity of a solid (σ) is described by an activated hopping processfrom one occupied lattice site to a neighboring empty latticesite (Figure 2a). This hopping process costs energy and ishence governed by a ΔG for the migration of the ion throughthe crystal structure. If the final and initial state are crystallo-graphically equivalent, a symmetric activation profile results,which leads to a high probability of a forward jump by the ionoccupying the transition state (i.e., the saddle point).35,36 Gen-erally, the Gibbs free energy of the hopping process is separatedinto an activation energy EA (i.e., the migration enthalpy) andthe entropy of migration ΔSm, leading to a conductivity of:

σσ

= −

Te E k T0 /A B

(1)

where the entropy of migration ΔSm is contained within theprefactor σ0. Using conventional hopping theory,37 theprefactor σ0 itself depends on a geometrical factor z thattakes into account different diffusion anisotropies and correla-tion factors, the density of charge carriers (i.e., the mobilespecies), the charge of the ions Ze, the entropy of migrationΔSm, the jump distance a0, and the jump frequency ν0 to afford:

σ ν= Δzn Zek

e a( ) S k

0

2

B

/0

20

m B

(2)

Meanwhile, the true measurable activation barrier is the sum ofthe enthalpy of migration ΔHm and the defect formationenthalpy ΔHf with

Figure 1. From the conventional battery architecture to solid-statebatteries with a lithium metal anode. Replacing the thin separator(gray band) and carbon anode (gray circles) with a solid electrolyte(orange circles) and lithium metal (light yellow), respectively, isexpected to increase the volumetric and gravimetric energy densities.However, the schematic already shows that the intimate grain contactrequired in a solid-state system necessitates sufficient percolation tofully access the cathode active material (violet circles), low microstruc-tural strain, high ionic conductivity, and good interfacial stabilities.Adapted with permission from reference 5. Copyright 2016 NaturePublishing Group.

Figure 2. (a) Schematic jump process of a moving ion in a solid. During the jump, the ion must bypass the transition state and overcome theassociated ΔG. (b) Hexagonal close-packed lattice with different diffusion pathways with (c) the calculated migration barriers. The face-sharingtetrahedra lead to a symmetric diffusion profile with a low activation barrier. Data in panel c were digitized from reference 73.

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= Δ + ΔE H H1/2A m f (3)

In superionic conductors, the defect formation enthalpy is oftennegligibly small, and the measured activation barriers corre-spond to the enthalpy of migration.31 However, the defect for-mation enthalpy may not always be neglected and can actuallybe used to estimate ionic conductivities.38,39

Certain material classes exhibit intrinsically high ionic con-ductivities for Li+ and Na+ ions, such as the NASICON family,40−44

the LISICON class,45−47 Li+-conducting garnets,48−51 as well as theLi+- and Na+-conducting thiophosphates within the Li10GeP2S12(LGPS)52−57 or Li6PS5X (argyrodite)26,27,58−61 families, amongother superionic thiophosphates.25,62−67 It should also be notedthat additional material classes can be found when alternativediffusing species are considered (e.g., α-AgI,32 RbAg4I5,

68,69 and(Ag/Cu)-argyrodites70,71). The common thread in all of thesematerials is that the underlying crystal structure is highlyfavorable for ionic transport. For instance, a large number ofavailable crystallographic sites, over which the moving ions aredistributed,34,38 is crucial.30−32 In other words, a high chargecarrier density is achieved if all ions can move and do not haveto migrate via an interstitial defect or vacancy. Furthermore, alow activation barrier for jumps between adjacent sites is alsonecessary (Figure 2), as well as similar potential energies at theinitial and final state of the jump. As previously mentioned,during the ion jump from the initial to final crystallographicstate, the mobile ion passes through a transition state or saddlepoint. Upon reaching the transition state, the ion may eitherundergo a forward jump to the final state or it may return backto the initial state. Ionic migration will only have occurred if theion proceeds forward to the final state. If the energy landscapesof the initial and final states are similar, the probability of aforward jump is high, thereby leading to a larger jump fre-quency, as the jump frequency is the product of the attemptfrequency and the jump probability.35,36 However, if the finalstate has a higher potential energy, the probability of a back-ward jump then becomes more likely. In terms of transition statetheory, the jump frequency can be described by the inverse of theexcited state lifetime, i.e. the lifetime of the transition state.72

Applying these concepts toward the design of SEs possessingsuitable energy landscapes, Ceder and coworkers recently showedthat body-centered cubic lattices can achieve a close proximityof lattice sites and favorable ion mobility with more symmetricenergy profiles from mostly face-sharing tetrahedra.73 Beyondenhancing the symmetry in the potential energy landscape,face-sharing polyhedra also lead to lower potential barriers,relative to edge-sharing polyhedra, as the open window for dif-fusion is geometrically wider and thus more favorable.73 Theinfluence of the structure on the migration barriers in a proto-typical hexagonal close-packed lattice is shown in Figures 2band c. The lowest migration barrier is found for the symmetricface-sharing tetrahedral jumps, whereas migration through anoctahedral site raises both the energy barrier and the probabilityof a backward jump.Knowing these static structural requirements (i.e., wide dif-

fusion pathways, broader bottlenecks for ionic jumps, and alarge carrier density) has aided in the design and optimizationof many SEs. Starting from α-AgI, Rb+ substitution was used inRbAg4I5 to stabilize the favorable cubic phase at room tempera-ture. Later, the Ag+- and Cu+-conducting cubic argyrodites werealso found. In the LISICON class of Li2+2xZn1−xGeO4, substi-tution for Ge with Ga, P, As, and V46,47,74−78 helped broadendiffusion pathways while simultaneously tuning the lithium

content. Here, changing the lithium content even alters the con-duction mechanism. For instance, the Li-rich compositionsexhibit an interstitial mechanism, while Zn-rich Li2+2xZn1−xGeO4undergoes vacancy-mediated conduction.45,47,78 Still today, theaforementioned SEs are being investigated for the transitionbetween defect-mediated hopping mechanisms and superionictransport, solely due to the diversity of polyanions.79 In a similarfashion, the nature of the polyhedral species has also been shownto affect ionic transport in the lithium thiophosphate glasses.25

The preparation of different solid solutions in NASICON-typeSEs (e.g., Li1+xAlxGe2−x(PO4)3 or Na1+xZr2P3−xSixO12), hasproven successful in increasing the content of the mobilespecies as well.32,41,80,81 Changing the lithium concentrationin thiophosphate-based classes such as the argyrodites andLi10GeP2S12 can also tailor the ionic conductivity.24,53,56,82−84

Even the best garnet-structured electrolyte Li7La3Zr2O12 wasobtained by increasing the lithium concentration from the ini-tially reported composition of Li5La3M2O12 (M = Nb, Ta).85

In addition to optimizing the carrier concentration, effortsoften focus more on broadening the diffusion pathways andopening up the geometric bottlenecks for ion jumps whileensuring that changes to the local arrangement of the neigh-boring ions does not disrupt the diffusion pathways.73,86,87 Thisapproach has been successful in the garnet-type ionic conduc-tors,88 the NASICON44,80,89 and LISICON classes,46,47,74−78 aswell as lithium and sodium ion conducting thiophos-phates.26,27,90,91 While being extensively employed throughoutthe literature, the general theme of “the broader the diffusionpathways, the lower the activation barriers” emerged.32 How-ever, some materials do not follow this trend. In Li10GeP2S12,isoelectronic substitutions with Si or Sn were used towardobtaining higher Li+ conductivities.56,92,93 While ab initio molec-ular dynamics simulations predicted that the Si compoundsshould exhibit higher conductivities,56 the chemical intuition oftargeting broader diffusion pathways suggested that the Sncompounds would in fact be better. To gain a better under-standing of the optimum channel size, Kato et al.94 investigateda series of Li10Ge1−xMxP2S12 (M = Si, Sn) solid solutions andconfirmed that the composition with the maximum conduc-tivity and the lowest, most favorable activation barrier is in factclose to the composition of Li10GeP2S12. Thus, despite the pre-dictions, elemental substitutions straying from the Li10GeP2S12composition led to lower ionic conductivities in the end. Figure 3aillustrates the lithium distribution in Li10GeP2S12 as obtainedfrom neutron diffraction data, showing the three-dimensionaldiffusion of Li+ along the tunnels of the z-axis and in the x−y-plane.57 In this structure, the smaller Si4+ leads to a lower con-ductivity, owing to a geometric restriction of the diffusion path-ways.94 However, the structural reasons behind the decrease inconductivity and the concurrent increase in the activationbarrier upon moving from Ge4+ to Sn4+ cannot be explained bythe ionic radii. As shown in Figure 3c, with increasing unit cellvolume and c/a ratio, the activation barrier indeed increasesand the ionic conductivity decreases.95 At the same time, adecrease in the S3−S3 distance with increasing unit cell volumecan also be observed (Figure 3d). This region of the structurethen acts as a bottleneck for ionic motion in the z-direction.The larger Sn4+ ions force these sulfur atoms closer together,which in turn causes the Li+ ions to jump through a narrowerwindow, destabilizing the transition state and raising theenergetic barrier for ionic motion (Figure 3b). It seems thatfor the composition of Li10GeP2S12, the interplay betweenpossessing large enough polyhedra and sufficiently broad

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bottlenecks for ionic jumps is optimal for transport. Thestructure−property correlations in Li10GeP2S12 demonstratethat while altering the unit cell volume and diffusion pathwaysis often a good first approach toward optimizing the transportin ionic conductors, more local structural changes in less-isotropic materials can also counter typical structural intuition.Additionally, it is also important to note that the in-planediffusion does not dominate the lithium mobility in this system,given that an increase in the c/a ratio does not result in anenhancement of the conductivity. Instead, the tunnels along thez-axis are the dominant transport pathways, and the in-planediffusion is likely only used to circumvent mobility issuesarising from point defects, as often observed in one-dimen-sional ionic conductors.96

The influence of static effects such as charge carrier densityand the breadth of the diffusion pathways has provided the fieldwith powerful tools for tuning the conductivity of ionicconductors. However, some structures do not allow for suchmodifications or even behave counterintuitively. Therefore, it isimportant to combine structural investigations with transportmeasurements, nuclear magnetic resonance, and theoreticalcalculations within the field of ionic conductors.Future directions will focus on gaining a better understand-

ing of local Coulombic interactions and possible inductive effectsin ionic conductors. For instance, in the Li10Ge1−xSnxP2S12 solidsolutions, the lower electronegativity and larger ionic radius of

Sn4+ leads to less electron density in the M4+−S2− bonds, ascompared to Ge4+.95 With increasing Sn content, the higherelectron density on the S2− atoms results in stronger Coulom-bic attractions between Li+ and S2−, thereby augmenting theactivation barrier and the prefactor for ionic motion (Figure 4).Similar inductive effects have been found in Na3P1−xAsxS4,

97

but a more in-depth theoretical approach using Bader or Born-effective charges is necessary to better understand inductiveeffects in solids. In addition to these more bonding- and structure-based approaches, a deeper understanding of correlated ionicmotion must also be achieved to better probe the motion of ionsin solids.98−100 Again, while often challenging, combined theoryand experimental investigations are required to shed light onhow one can initiate and tailor correlated motions, i.e. identifythe design principles that activate such correlated motion effects,which may be beneficial for ionic transport.

2.2. Dynamic Lattice and Prefactor Dilemma. In addi-tion to the static influences from the underlying crystal struc-ture, increasing the polarizability of the anion framework hasalso been suggested to lower the activation barriers for ionjumps.73 Considering Figure 2a, a more polarizable sublatticemeans softer bonding interactions and a lower energetic costfor the displacement of the lattice during a jump. The idea thatthe “softness” and dynamics of a host lattice have an influenceon the ionic motion in solids had initially been explored in the1970s/80s,101−105 as the ionic jump processes are thermally

Figure 3. (a) Reconstructed lithium negative nuclear density maps of Li10GeP2S12, illustrating the diffusion along the tunnels in the z-direction andthe x−y-plane. (b) Li−Li jump in the z-direction with the S3−S3 distance as the geometric bottleneck. (c) Increasing activation barrier despitebroader diffusion pathways in the x−y-plane and larger c/a ratio in Li10Ge1−xSnxP2S12. (d) Interdependence of the bottleneck S3−S3 distance on theconductivity and activation barriers. Figure in panel a is adapted with permission from reference 57. Copyright 2016 American Chemical Society.Figures in panels b−d are adapted with permission from reference 95. Copyright 2018 American Chemical Society.

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activated and must therefore be closely related to the phononspectrum.106 As the phonon spectrum and its collective vibra-tional motions directly correlate to the elasticity of the lattice,the covalency and polarizability affect the oscillator strength ofthe lattice and with it the transition rate of the hoppingion.105,107,108 Shao-Horn et al. recently highlighted the influ-ence of the phonon frequency (and therefore phonon bandcenters in the phonon dispersion curve) on the activationbarrier, showing that a softer lattice and lower phonon frequen-cies were correlated with a decrease in the associated activationbarriers.109 The expected influence of a softer, more polarizablelattice on the ionic transport is depicted in Figure 5a, in whicha softer lattice should lower the activation barrier and theeigenfrequency of the oscillations, i.e. the oscillator or attemptfrequency, through a broadening of the local jump oscillators.Beyond the strength of the lattice vibrations, concerted polyhe-dral rotations have also been shown to assist the mobile ionthrough a so-called paddle wheel mechanism.110,111

The influence of lattice softening was recently corroboratedexperimentally using solid solutions of Li6PS5X (X = Cl, Br, I)for Li+ ions and Na3PS4−xSex for Na

+ ions, in which the anionpolarizability was systematically varied.26,90 In the absence oflocal structural changes, the activation barrier for ionic motiondecreases with decreasing Debye frequency of the lattice (υD),which directly relates to the softness of both the phononspectrum and the lattice.26 Figure 5b shows the dependence ofthe activation energy on a softening lattice in Na3PS4−xSex. Inaddition to affecting EA, the prefactor for ionic motion alsodecreases over orders of magnitude with a softer lattice. Consid-ering eq 2, the charge carrier density is not affecting the trans-port, given the isoelectronic nature of the substitutions, andfurther, changes in the attempt frequency and jump distance arealso not large enough to account for the significant decrease ofthe prefactor.26,90 This then leaves the entropy of migration as themain culprit for the detrimental influence of a soft lattice on thetransport, despite the concurrent decrease in the activation energy.The entropy of migration itself depends on the phononic

properties of a material and can, for a small vibrational approxi-mation, be expressed as112−114

ν

νΔ =

Π

Π

=

−

=

−

⎛

⎝

⎜⎜⎜

⎞

⎠

⎟⎟⎟S k ln i

N

i

Nm B1

3 1

iI

1

3 1

iS

(4)

with the normal frequencies ν1...νN referring to vibrationsaround the initial state (I) and the saddle point or transitionstate (S). While the entropy of migration is affected by nearlyall vibrational frequencies around the diffusional pathway, themigration enthalpy is related to the one vibrational mode thatcarries the ion across the saddle point.47 This attempt fre-quency is typically approximated by the Debye frequency.26

Therefore, a softening of the lattice leads to lower Debyefrequencies (Figure 5b) and is expected to not only affect theactivation barrier but also the prefactor through a change in boththe oscillator frequency and the entropy of migration.26,90,115

Interestingly, the decreasing prefactor with increasing latticesoftness leads to a dilemma, as the approach of introducingmore polarizable anions has always been thought to be bene-ficial for ionic transport. Hence, the influence of the phononfrequencies on the entropy of migration has been largely over-looked and with a decreasing activation barrier, a decrease inthe prefactor is also to be expected (Figure 5c). This is theso-called Meyer−Neldel rule,116−120 in which the slope of theMeyer−Neldel plot has been linked to the entropy of migrationand the Debye frequencies of the lattice.121 A stiffer lattice maythen result in a flatter slope in the Meyer−Neldel plot, while asofter lattice would exhibit a steeper slope. Indeed, for thestiffer Li10Ge1−xSnxP2S12 compounds,

95 the slope seems to beflatter relative to the softer Li6PS5X (X = Cl, Br) andNa3PS4−xSex compounds. Thus, materials possessing a flat slopein the Meyer−Neldel plot (i.e., stiffer) are not expected tonegatively impact the prefactor during the optimization of theactivation barrier. A notion that is counterintuitive to theknown paradigm of “the softer the lattice, the better”. It maythen be beneficial to start with stiffer structures and subse-quently soften the lattice, instead of starting with soft materialsdirectly, when optimizing ionic transport. On the other hand, insoft materials with a steep slope, chemical modifications towardincreasing the prefactor, such as increased charge carrier den-sities, may then be less detrimental to the activation barrier.Still, when comparing classes of materials, the more polarizableanion lattices are highly favorable. Moreover, the Debye fre-quencies of SEs vary with composition,26,90,95 and the Meyer−Neldel slope may be decreasing with increasing Debye fre-quency (i.e., at higher activation barriers). All of the above-mentioned considerations will likely be helpful for the optimi-zation of transport in SEs. Accordingly, an alternative approachcould be to find structures with soft vibrational modes in the

Figure 4. (a) Activation energy EA and the prefactor σ0 show an increase with increasing Sn content, corresponding to a decrease in the M4+−S2−interactions. (b) Sketch illustrating the altered Coulombic interactions between sulfur and lithium, resulting from the different M−S-bondingcharacteristics. Figure is adapted with permission from reference 95. Copyright 2018 American Chemical Society.

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excited state and stiffer phonon modes in the initial groundstate to simultaneously obtain high prefactors and low enthalpiesof migration.Both examples, involving the Li6PS5X and Na3PS4−xSex

systems, show that the ambiguous influence of lattice dynamicsand the “prefactor dilemma” have been overlooked in the pastand that a better understanding of these concepts in relation toionic transport is required. Further still, recent studies on theinfluence of a dynamic lattice on the ionic transport havealready shown that the entropy of migration and the prefactor

cannot be ignored.122 In particular, when using theoreticalapproaches to compute ionic conductivities, as well as the datamining for possible structures with high ionic conductivities, itis necessary to include entropic considerations in addition tothe calculations of activation barriers.29

2.3. Average and Local Structures. While typicalapproaches toward enhancing ionic transport involve tailoringthe crystal structure through compositional modifications, itis not always clear if the materials are also being influencedby secondary amorphous phases. In particular, in the class ofthiophosphate ionic conductors, the glasses and glass ceramicsare receiving a lot of attention for this reason.66,67,123−127 Thepoint is that investigations into material behavior using onlyBragg diffraction are proving to be insufficient. Quite often,Raman spectroscopy and nuclear magnetic resonance also needto be used to resolve the different species contributing to thematerial properties.115,128,129 Dietrich et al. even used pairdistribution function (PDF) analysis to show that lithiumthiophosphates, which appear crystalline by Bragg diffraction,may actually be glass-ceramics.38 These secondary phases,despite revealing no coherent scattering domains,130 may notbe entirely amorphous and could be significantly influencingthe ionic transport. In poorly conducting crystalline materials,such glassy phases may be acting as conducting fillers,38 whereasin good conducting materials, these phases may actually hinderthe ionic transport.131 The synthetic conditions must thereforebe optimized to attain the appropriate balance of glassy andcrystalline phases for good transport behavior.Importantly, the synthetic conditions can also influence the

resultant crystal structures in thiophosphate ionic conductors.In many syntheses, mechanical alloying, i.e. ball milling, is com-monly used. For example, Li10GeP2S12 can be achieved onlythrough such milling techniques. On the other hand, Na3PS4crystallizes in a tetragonal structure when prepared via classicalhigh-temperature routes,132 but mechanical alloying leads to thecubic polymorph.11,133 Here, the cubic polymorph exhibitsionic conductivities orders of magnitude better than the tetrag-onal polymorph, despite theoretical predictions suggesting thatboth should possess similar conductivities.134−136 Using PDFanalysis, Krauskopf et al. were able to show that the localstructures of both the tetragonal and cubic polymorphs areactually tetragonal (Figure 6). In other words, both polymorphsshow the same tetragonal structural motif on the local scale,even though the average crystal structure suggests differences.28

While the typically performed Bragg diffraction gives only aglobally averaged depiction of the structure, analysis of the pairdistribution function provides a local structural picture, as itrepresents a histogram of atom−atom distances within the solidin real space. Notably, upon rapidly annealing the high-performancecubic polymorph, a tetragonal structure can be obtained, asindicated by Bragg diffraction; however, the transport proper-ties remain unchanged.28 It is already known that in materialspossessing, for local dipoles (e.g., ferroelectrics), there can bestructural differences between the local and average lengthscales,137−139 and it appears that this is also possible for ionicconductors as well.The differences in the local structure, the enhanced transport

arising from processing techniques, and the influence of under-lying glassy phases demonstrate that the structures, composi-tions, and properties of the thiophosphate SEs are all interre-lated in a complex manner. In the superionic thiophosphates,the transport behavior does not necessarily depend on thecrystal structure alone, but rather the microstructure, the

Figure 5. (a) Schematic of the effect of lattice softening on the ionictransport. With increasing softness of the lattice, the local vibrationalfrequencies decrease along with the activation barrier for the jump.(b) Decreasing activation barrier EA and prefactor σ0 with decreasingDebye frequency νD in Na3PS4−xSex. (c) Meyer−Neldel plot showingthe interdependence between the prefactor and activation barrier.Figure in panel a is adapted with permission from reference 26.Copyright 2017 American Chemical Society. Data in panels b and cwere taken from references 26, 90, and 95.

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multifaceted nature of the glass-ceramic, or even defects inducedby harsh synthetic conditions. Ultimately, more work mustbe done to answer the many open questions concerning thestructure−property relationships at play in these complicatedsystems.

3. INTERFACIAL PROCESSES IN SOLID-STATEBATTERIES EMPLOYING THIOPHOSPHATES

The importance of deciphering the principles that govern ionicconduction in solids is clear, given that achieving higher energydensity batteries will be possible only through the use of thickelectrode configurations coupled with thin separator layerspossessing exceptional ionic conductivities.17 Importantly, suchthick electrode configurations in SSBs may necessitate evenhigher ionic conductivities to prevent additional cell over-potentials.17 Nevertheless, the ionic conductivity of the SE isnot the only bottleneck of the performance for SSBs. Duringbattery operation, reactions between the electrode materialsand the electrolyte evolve resistive interfacial layers thatstrongly affect the subsequent capacity and cycle life. Thus, itis imperative that we dig deeper into both the identity of the

interfacial species as well as their influence on the deviceproperties.

3.1. Interfacial (In)stabilities. In viewing the schematic ofan archetypical SSB (Figure 1), one can already glean thesignificant role that interfaces play in influencing the batteryperformance.140 Importantly, degradation mechanisms of solidelectrolytes have been reported throughout the literatureregarding both electrode interfaces. On the anode side, thelithium metal electrode reduces many of the best-performingSEs, producing resistive interlayers.3,4,12−15 For example, Wenzelet al. monitored the decomposition of LGPS in contact with Limetal using in situ X-ray photoemission spectroscopy (XPS)and electrochemical impedance spectroscopy. Therein, thecontinuous growth of an interphase composed of Li2S, reducedGe species, and Li3P was observed, with a correspondingincrease in the interfacial resistance.14 While the reaction of Limetal with various metal-containing solid electrolytes leads tothermodynamically unstable, mixed-conducting interphases thatfacilitate continuous SE degradation, compositional tailoring ofSEs may enable the realization of more favorable, kineticallystable interphases. In other words, one potential approach for

Figure 6. Structural representations of the cubic (a) and tetragonal (b) crystal structures, showing the differences in local Na+ ordering. Fitting of theexperimentally obtained G(r) data for the cubic polymorph (c-Na3PS4) using both the cubic (c) and tetragonal (d) structural models, therebycorroborating that the tetragonal motif is the prevalent structural motif on the local scale. Figure is reproduced with permission from reference 28.Copyright 2018 American Chemical Society.

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mitigating SE degradation could involve the identification of SEcompositions that only partially react with Li metal to formthin, ionically conducting, yet electronically insulating inter-phases.14,141 Additional concerns regarding the anode includedendrite and crack formation, thus making it necessary to findways to optimize critical current densities.142 Thin film coatingson lithium metal are also being developed to hinder degrada-tion and improve the wetting behavior at the anode interface.143

Meanwhile, on the cathode side, oxidative degradation of thio-phosphate SEs against high-voltage active materials is expected4,13

and has also been experimentally corroborated.19,144−151 Thesereports show that the oxidative SE degradation during cyclingstems from the intrinsically narrow electrochemical stabilitywindow of the electrolytes. A hypothesis that was furthercorroborated by first principles calculations from Mo andcoworkers.3,4,152 Interfacial reactions can also result in transition-metal migration from the cathode active material, hindering thebattery performance even more.145 In the end, interfacial reac-tions involving thiophosphates generate ionically insulatingproducts such as Li4P2S7, Li4P2S6, and Li2P2S6.

153,154 Therefore,the development of protective coatings for cathode activematerials is of paramount importance. Zhang et al. recentlyextended both the rate capability and the capacity retention incells constructed with the cathode active material LiCoO2 byemploying a LiNb0.5Ta0.5O3 coating.16 The presence of theamorphous oxide layer lowered the interfacial resistance whilealso likely inhibiting deleterious interfacial reactions bymitigating direct contact between the cathode material andthe SE. Interestingly, these interfacial reactions may not neces-sarily be irreversible. Redox-active behavior of lithium thio-phosphates has already been reported in the literature,151,155−158

though an investigation into the active contribution of the SE tothe cell performance is still missing.To better understand the instability of thiophosphate SEs

against high-voltage cathode materials, Koerver et al. used elec-trochemical impedance spectroscopy and in situ X-ray photo-emission spectroscopy to demonstrate the redox-activity ofβ-Li3PS4 induced upon battery operation.18 Notably, thesereactions are strongly dependent on the state of charge of thebattery, as well as the applied upper cutoff potential. During thecharging of the cell, the oxidation of the SE in contact withLiNi0.8Co0.1Mn0.1O2 (NCM-811) can be seen through thedevelopment of an additional interfacial resistance, relative tothe impedance in the discharged state (Figure 7a). With thehelp of in situ XPS, the evolution of the chemical species couldbe resolved, thereby explaining the interfacial phenomena

(Figures 7b−d). During the degradation of β-Li3PS4, the oxi-dized products polymerize into interconnected P−[S]n−P-typenetworks, which may lead to the gradual formation of uniquelyconducting phases.147,151,156,159 Ultimately, SE degradation atthe cathode results in decreased capacity retention through thedevelopment of a semi-irreversible cathode/SE interfacialcharge transfer resistance, further highlighting the need forsuitable protective coatings for the cathode active material. Allof the above-mentioned conditions occurring at the anode andcathode drive the search for stable (or self-limiting) SE inter-phases to enable long-term stability and the use of lithiummetal anodes.

3.2. Microstructural and Volumetric Influences. Toelectrochemically address all of the active particles simulta-neously and avoid local overcharging, achieving adequate perco-lation and contact within the electrodes is crucial. Solid-statebatteries present a different situation from liquid electrolyte-based lithium-ion batteries, where pore diffusion of the liquidelectrolyte provides sufficient percolation.5 Thus, added con-siderations (e.g., particle size/morphology of the active materialand composite electrode composition) must be taken intoaccount when optimizing SSB configurations.16,160 Accordingly,Janek and coworkers were able to show that the ratio of SE tocathode active material in the cathode composite is indeedinfluential.16 While a high loading of active material leads tohigher capacities, the concurrently lower volume fraction of theSE decreases the attainable current densities. It seems to benecessary to design the ratios in the electrode based on thedesired application, while also keeping the composite tortuosityin mind.17 For example, thicker electrodes with a higher ratio ofcathode active material are required for high-energy applica-tions, while obtaining higher power may require faster ionicconductors or a larger amount of the SE.16,17

Furthermore, good electrode percolation has also been achievedthrough slurry-based electrode preparations.161−163 More com-monly, maintaining sufficient electrode contact in SSBs usingthiophosphate-based electrolytes is accomplished through theapplication of external pressures. Beyond the initial contact, themechanically soft thiophosphate electrolytes (Young’s modulusof approximately 20 GPa22,164,165) will also undergo micro-structural deformations during the cycling-induced expansionand contraction of cathode active materials (e.g., LiCoO2 andLiNi1−x−yCoxMnyO2). Moreover, it was recently shown thatdespite the soft nature of the thiophosphates, which may allowfor the correction of elastic mismatches during cycling, a low frac-ture toughness exists that may generate additional microstructural

Figure 7. (a) Evolution of the interfacial resistance between the cathode and lithium thiophosphate SE. The resistance is heavily dependent on thestate of charge of the battery due to the redox active behavior that can be monitored using in situ X-ray photoemission spectroscopy, shown in panelsb−d. Figures reproduced with permission from reference 18. Copyright 2017 Royal Society of Chemistry.

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flaws.10 Meanwhile, in cathode active materials, the volumechanges are structurally related to the deintercalation, theresultant Coulombic repulsion between the MO2 layers, andthe changing ionic radii of the transition metals.166,167 The crys-tallographic volume changes occurring within LiCoO2 andNCM-811 during deintercalation are provided in Figure 8a.During charging (i.e., delithiation), the overall volume of theLiCoO2 unit cell expands, while the unit cell volume of NCM-811 contracts. It should also be mentioned that similar consid-erations must also be made for the anode side, as most anodematerials will experience volume expansion during lithiuminsertion.20,168 While liquid electrolytes can compensate thesevolume changes, a tremendous microstructural strain will buildup along the grain contacts in SEs.20 Figure 8b shows thenominal strain response upon cycling in an SSB using a metalanode and LiCoO2 as the cathode active material. During charg-ing, the volumes of both the anode and the cathode expand,leading to a reversible pressure increase.20 The resulting pres-sure change of approximately 1 MPa was found to promotecracking and bending within the SSB due to local microstruc-tural strain.20

When considering volumetric changes, it becomes clear thatthe pressure inside the cell will vary accordingly. As previouslymentioned, for a material like NCM-811, the composite cath-ode exhibits an overall volume contraction upon delithiation.However, given the reversibility of these volume and pressurechanges, one must also consider the microstructural implica-tions within solid-state cells. Figure 8c shows a colorized scanning

electron micrograph of NCM-811 after cycling in a solid-statebattery. Due to the volume contraction during charge, thespherical particles have lost the intimate contact that they ini-tially had with the SE.19 This contact loss leads to an additionalinterfacial resistance and prevents the full discharge (lithiation)of the material, thereby contributing to the typically observed,larger capacity loss found in the first cycle when using NCM-811in solid-state batteries,169,170 as compared to liquid electrolyte-based cells (Figure 7d).19

The influence of electrochemically induced local strain mustbe considered when designing SSBs, especially if they are to berun without external pressure and as such, a certain flexibilityneeds to be provided166,167 to mitigate volume changes duringcycling. It has already been computed that recovering capacitysolely by increasing external pressure, thereby reversing contactloss, would not be feasible.171 Therefore, strategies for reducingnegative volumetric effects, without employing added externalpressure, will be extremely important moving forward.

4. SUMMARY AND FUTURE DIRECTIONS

In this perspective, we discussed the structural influences on theionic conduction in solids as well as the interfacial chemistry insolid-state batteries. We highlighted the notion that the typicalchemical intuition of attaining broader diffusion pathways maynot always be correct and that more in-depth structural inves-tigations are necessary to identify potential bottlenecks for ionicdiffusion. We showed that there is a prefactor dilemma whentrying to utilize softer, more polarizable lattices, given the

Figure 8. (a) Change in unit cell volume during charge (delithiation) of LiCoO2 and NCM-811. (b) Nominal cell pressure during cycling of an SSBusing an In anode and LiCoO2 cathode composite. (c) False color scanning electron micrograph showing the contact loss between NCM-811 andthe SE. The volume contraction during the first charge of NCM-811 leads to a lower first cycle efficiency in the solid-state cell, as compared to a cellusing a liquid electrolyte, shown in panel d. The purple data correspond to the SSB cell, and the orange to the liquid electrolyte cell. Data for panel aare extrapolated from references 166 and 167. Panel b is reproduced with permission from reference 20. Copyright 2017 Royal Society of Chemistry.Panels c and d are reproduced with permission from reference 19. Copyright 2017 American Chemical Society.

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interdependence between the activation barrier and the pre-factor. These results show that the dynamics of the lattice arean important factor governing ionic conduction and that moreresearch is necessary to better understand the associated influ-ences and to find better descriptors for ionic motion in solids.Additionally, we showed that compositional modifications canalter the electrostatic interactions within the lattice, causinginductive effects, which also need to be considered andexplored in ionic conductors. We further discussed how thelocal structures may also differ from the average structures insolid electrolytes, depending on the synthetic conditions, andhow this correlates to the transport behavior.Regarding the interfacial processes occurring in SSBs, we

showed that the interfaces and interphases between theelectrodes and the SEs must be designed to provide long-term performance and stability. Whether the approach involvesprotective layers at the anode and/or thin coatings of high-voltage cathode materials, decomposition of the SE separatorsneeds to be mitigated if not entirely inhibited. Moreover, wedemonstrated that the volumetric changes taking place withinthe electrodes lead to severe and often detrimental micro-structural effects that further limit SSB performance. Thus, theoptimization of SSB architectures will require the inclusion ofboth interfacial and microstructural considerations to advancethis technology.We hope that this perspective enables future strategies toward

understanding ionic conduction in solids and the optimization ofsolid-state batteries. In the end, the underlying principles gov-erning ionic motion, in addition to the interfacial reactionsoccurring between the electrodes and the solid electrolyte, mustbe clarified to evaluate and improve solid-state battery technolo-gies in a more realistic fashion.

■ AUTHOR INFORMATIONCorresponding Author*E-mail: wolfgang.g.zeier@phys.chemie.uni-giessen.de.ORCIDWolfgang G. Zeier: 0000-0001-7749-5089NotesThe authors declare no competing financial interest.Biographies

Sean Culver received his Ph.D. in Inorganic Chemistry in 2016 fromthe University of Southern California in Los Angeles under thesupervision of Prof. Richard Brutchey. He is currently an Alexandervon Humboldt postdoctoral fellow at the Justus-Liebig-UniversityGiessen, working with Prof. Jurgen Janek and Dr. Wolfgang Zeier. Hisresearch interests include fundamental structure−property relation-ships in ionic conductors in addition to the interfacial chemistry andoptimization of all-solid-state batteries.

Raimund Koerver received his M.Sc. in Chemistry from the Ruhr-University Bochum. During his Master’s thesis he was a visitingresearch fellow at the Monash University, Melbourne. Currently, he isa Ph.D. candidate at Justus-Liebig-University Giessen under super-vision of Prof. Jurgen Janek and Dr. Wolfgang Zeier. His researchinterests include interfacial reactions and (chemo-)mechanical effectsin thiophosphate-based all-solid-state batteries.

Thorben Krauskopf received his M.Sc in Chemistry from the Justus-Liebig-University Giessen. Currently, he is a Ph.D. student at Justus-Liebig University Giessen under the supervision of Prof. Jurgen Janekand Dr. Wolfgang Zeier. His research interest includes thefundamental structure−property relationships in ionic conductors

and the transfer kinetics between metal anodes and alkali solidelectrolytes.

Wolfgang Zeier received his Ph.D. in Inorganic Chemistry in 2013from the Johannes-Gutenberg University in Mainz under thesupervision of Prof. Wolfgang Tremel and Prof. Jeffrey Snyder(California Institute of Technology). After postdoctoral stays at theUniversity of Southern California and at the California Institute ofTechnology, he was appointed group leader at the Justus-Liebig-University Giessen, within the framework of an Emmy-Noetherresearch group. His research interests encompass the fundamentalstructure−property relationships in solids with a focus on thermo-electric and ion-conducting materials as well as solid−solid interfacialchemistry in all-solid-state batteries.

■ ACKNOWLEDGMENTS

The research was supported by the Deutsche Forschungsge-meinschaft (DFG) under Grant ZE 1010/4-1. S.C. gratefullyacknowledges the Alexander von Humboldt Foundation forfinancial support through a Postdoctoral Fellowship. R.K.gratefully acknowledges financial support by the Funds of theChemical Industry (FCI).

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Chemistry of Materials Perspective

DOI: 10.1021/acs.chemmater.8b01293Chem. Mater. 2018, 30, 4179−4192

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