Designing ionic conductors: the interplay between structural

Subscriber access provided by University | of Minnesota Libraries

Perspective

Designing ionic conductors: the interplay between structural phenomena and interfaces in thiophosphate-based solid-state batteries Sean P. Culver, Raimund Koerver, Thorben Krauskopf, and Wolfgang G. Zeier Chem. Mater., Just Accepted Manuscript • DOI: 10.1021/acs.chemmater.8b01293 • Publication Date (Web): 11 May 2018 Downloaded from http://pubs.acs.org on May 13, 2018

Just Accepted “Just Accepted” manuscripts have been peer-reviewed and accepted for publication. They are posted online prior to technical editing, formatting for publication and author proofing. The American Chemical Society provides “Just Accepted” as a service to the research community to expedite the dissemination of scientific material as soon as possible after acceptance. “Just Accepted” manuscripts appear in full in PDF format accompanied by an HTML abstract. “Just Accepted” manuscripts have been fully peer reviewed, but should not be considered the official version of record. They are citable by the Digital Object Identifier (DOI®). “Just Accepted” is an optional service offered to authors. Therefore, the “Just Accepted” Web site may not include all articles that will be published in the journal. After a manuscript is technically edited and formatted, it will be removed from the “Just Accepted” Web site and published as an ASAP article. Note that technical editing may introduce minor changes to the manuscript text and/or graphics which could affect content, and all legal disclaimers and ethical guidelines that apply to the journal pertain. ACS cannot be held responsible for errors or consequences arising from the use of information contained in these “Just Accepted” manuscripts.

is published by the American Chemical Society. 1155 Sixteenth Street N.W., Washington, DC 20036 Published by American Chemical Society. Copyright © American Chemical Society. However, no copyright claim is made to original U.S. Government works, or works produced by employees of any Commonwealth realm Crown government in the course of their duties.

Page 1 of 35 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Chemistry of Materials

Designing ionic conductors: the interplay between structural phenomena and interfaces in thiophosphatebased solid-state batteries Sean P. Culvera, Raimund Koervera, Thorben Krauskopfa, Wolfgang G. Zeier*a a

Institute of Physical Chemistry, Justus-Liebig-University Giessen, HeinrichBuff-Ring 17, D-35392 Giessen, Germany.

Abstract Elucidating the underlying structural principles that govern ionic transport in thiophosphate solid electrolytes will enable the discovery of novel ionic conductors. Additionally, improving the properties of ionic conductors and exacting control over interfacial reactions and interphase stabilities are critical to the advancement of solid-state batteries. In this perspective, we focus on two major aspects at the foundation of solid-state battery development. First, we address the typical static structural requirements for achieving high ionic conductivities within thiophosphates, which is then extended to how a dynamic lattice and local structural effects can influence ionic transport. Furthermore, we provide an overview of some of the challenges that are currently hindering the progress of solid-state battery research, with particular attention being paid to interfacial instabilities and mechanochemical effects. We hope that this perspective provides a unique outlook on ionic conduction in thiophosphates toward the design of future solid electrolytes and highlights the importance of interfacial chemistry in the optimization of solid-state battery devices.

1 ACS Paragon Plus Environment

Chemistry of Materials 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

1. Introduction. Lithium-ion batteries are a ubiquitous energy storage technology that has transformed the role of commercial electronics in society. The high energy and power densities, along with the good reliability and cyclability have made lithium-ion batteries an obvious choice for powering our commercial electronic devices.1 While today’s state-of-the-art batteries offer good performances, further improvements to the energy density are not expected without the use of a lithium metal anode and high-voltage cathodes.2 Unfortunately, the commonly used liquid electrolytes are easily oxidized at higher voltages and are not able to suppress dendrite formation when lithium metal anodes are employed, leading to a variety of safety concerns.3,4

Figure 1: From the conventional battery architecture to solid-state batteries with a lithium metal anode. Replacing the thin separator (grey band) and carbon anode (grey circles) with a solid electrolyte (orange circles) and lithium metal (light yellow), respectively, is expected to increase the volumetric and gravimetric energy densities. However, the schematic already shows that the intimate grain contact required in a solid-state system necessitates sufficient percolation to fully access the cathode active material (violet circles), low microstructural strain, high ionic conductivity and good interfacial stabilities. Adapted with permission from reference 5. Copyright 2016 Nature Publishing Group. However, the use of a solid electrolyte (SEs) separator may circumvent the aforementioned problems. Recent efforts in SE development have provided the field with a variety of exceptionally 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 separator should mitigate unwanted electrode cross-talk of polysulfides or transition metal cations,8 while also eliminating typical leakage issues associated with liquid-based architectures. (2) A negligible partial electronic conductivity should prevent self-discharge. 2 ACS Paragon Plus Environment

Page 2 of 35

Page 3 of 35 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


Moreover, (3) the ionic transference number of nearly unity means only lithium ions are moving, so there should not be any polarization resistance at high current densities.7 (4) Chemical stability at elevated temperatures should also be improved, while additionally preventing the “freezing out” of the electrolyte at low temperatures. Finally, (5) the superior mechanical rigidity may even prevent dendrite formation caused during the electro-deposition of lithium.9,10 While dendrite formation still remains a challenge at high current densities and a variety of electrolyte-based interfacial instabilities have been reported against metal anodes,3,4,11–15 the use of a lithium metal anode will 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 anode is illustrated in Figure 1. While the decreased cell volume and thin metal anode provide gains in gravimetric and volumetric energy densities, the schematic already alludes to potential design issues. For example, the separator needs to be exceptionally thin, while at the same time preventing short-circuits. Moreover, ideal percolation must be achieved to electrochemically access all of the active material16 and a high ionic conductivity in the electrolyte is required to reduce overpotentials when increasing electrode thicknesses.17 Persistent grain contacts may also lead to interfacial reactions, high charge transfer resistances and microstructural strain.18–20 Ultimately, all of the aforementioned concerns have led to the predominant usage of high-performance thiophosphates21–23 in SSB research. Nevertheless, despite strong contributions from thiophosphate-based ionic conductors,6,7,24 the list of materials exhibiting sufficient ionic conductivity for practical SSB applications remains exceedingly short. While a multitude of approaches for enhancing ionic conductivity have been developed over the years (e.g. tuning of the crystal structure,25 elemental substitutions,26,27 microstructural modifications28), the success of any given approach is strongly dependent on our understanding of ionic conduction in solids.29 Thus far, much is already known regarding static structural influences on transport behavior, however, recent studies have pointed toward dynamic and local structural effects as well.30–32 Therefore, by expanding our knowledge of the complex processes governing ionic diffusion in thiophosphates, we can further enhance relevant transport properties and develop more general strategies for effectively tuning ionic mobility in SEs.



Page 4 of 35

In this perspective, we provide an overview of contemporary approaches for understanding and designing lithium thiophosphate solid electrolytes. First, we focus on the typical structural requirements for good ionic conduction, while showing that structural changes must be monitored in order to elucidate the conduction mechanisms and effectively tailor the performance of thiophosphates. We then present some of the recent progress on determining the influence of lattice dynamics on ionic transport, and further, we discuss the implications associated with disagreeing local and average structures within ionic conductors. Finally, we address the current challenges in SSBs concerning interfacial instabilities, redox activity of the electrolytes and mechanochemical changes that occur during cycling. We hope that by better defining the structure-property relationships at play in solid ionic conductors, in addition to clarifying the interfacial reactions and performance limitations that are currently impeding battery development, we can guide future efforts in the fields of ionic conductors and solid-state batteries.

2. Structural influences on ionic conduction Solid ionic conductors have been studied for decades and as such, have become indispensable for a myriad of applications.1,5,32–34 In particular, these ionic conductors have been heavily investigated as solid electrolytes (i.e. separators) in solid-state batteries. However, to be truly applicable, the target materials must possess exceptional transport properties. To accomplish this goal, improvements on the existing design principles of SEs, from a structural standpoint, need to be made. In considering how to approach this complicated task, one must develop a strong understanding over the structural influences governing ionic conduction in solids, thereby defining the appropriate routes to success. 2.1 Static lattice effects In general, the ionic conductivity of a solid () is described by an activated hopping process from one occupied lattice site to a neighboring empty lattice site (Figure 2a). This hopping

process costs energy and is hence governed by a ∆ for the migration of the ion through the

crystal structure. If the final and initial state are crystallographically equivalent, a symmetric activation profile results, which leads to a high probability of a forward jump by the ion occupying the transition state (i.e. the saddle point).35,36 Generally, the Gibbs free energy of


Page 5 of 35 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


the hopping process is separated into an activation energy (i.e. the migration enthalpy) and

the entropy of migration ∆ , leading to a conductivity of:

=

∙

Eq. (1)

where the entropy of migration ∆ is contained within the pre-factor . Using conventional

hopping theory,37 the pre-factor itself depends on a geometrical factor z that takes into account different diffusion anisotropies and correlation factors, the density of charge carriers

(i.e. the mobile species), the charge of the ions Ze, the entropy of migration ∆ , the jump

distance and the jump frequency to afford:

=

² ∆ ! ·

Eq. (2)

Meanwhile, the true measurable activation barrier is the sum of the enthalpy of migration

∆" and the defect formation enthalpy ∆"# with:

= ∆" + 1'2 ∆"#

Eq. (3)

In superionic conductors, the defect formation enthalpy is often negligibly small and the measured activation barriers correspond to the enthalpy of migration.31 However, the defect formation enthalpy may not always be neglected and can actually be used to estimate ionic conductivities.38,39 Certain material classes exhibit intrinsically high ionic conductivities 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 the Li+ - and Na+ - conducting thiophosphates within the Li10GeP2S12 (LGPS)52–57 or Li6PS5X (Argyrodite)26,27,58–61 families, among other superionic thiophosphates25,62–67. It should also be noted that, additional material classes can be found when alternative diffusing species are considered (e.g. α-AgI,32 RbAg4I568,69 and (Ag/Cu)-argyrodites70,71). The common thread in all of these materials is that the underlying crystal structure is highly favorable for ionic transport. For instance, a large number of available crystallographic sites, over which the moving ions are distributed,34,38 is crucial.30–32 In other words, a high charge carrier density is achieved if all ions can move and do not have to migrate via an interstitial defect or vacancy. Furthermore, a low activation barrier for jumps between adjacent sites is also necessary (Figure 2), as well as similar potential energies at the initial and final state of the jump. As previously mentioned, during the ion jump from the initial to final crystallographic state, the mobile ion passes through a transition state or saddle point. Upon reaching the 5 ACS Paragon Plus Environment


Page 6 of 35

transition state, the ion may either undergo a forward jump to the final state or it may return back to the initial state. Ionic migration will only have occurred if the ion proceeds forward to the final state. If the energy landscapes of the initial and final states are similar, the probability of a forward jump is high, thereby leading to a larger jump frequency, as the jump frequency is the product of the attempt frequency and the jump probability.35,36 However, if the final state has a higher potential energy, the probability of a backward jump then becomes more likely. In terms of transition state theory, the jump frequency can be described by the inverse of the excited state lifetime, i.e. the lifetime of the transition state.72 Applying these concepts toward the design of SEs possessing suitable energy landscapes, Ceder and coworkers recently showed that body-centered cubic lattices can achieve a close proximity of lattice sites and favorable ion mobility with more symmetric energy profiles from mostly face-sharing tetrahedra.73 Beyond enhancing 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 diffusion is geometrically wider and thus, more favorable.73 The influence of the structure on the migration barriers in a prototypical hexagonal close-packed lattice is shown in Figure 2b and 2c. The lowest migration barrier is found for the symmetric face-sharing tetrahedral jumps, whereas migration through an octahedral site raises both the energy barrier and the probability of a backward jump.

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 the associated ∆. b) Hexagonal close-packed lattice with different diffusion pathways with c) the calculated migration barriers. The face-

sharing tetrahedra lead to a symmetric diffusion profile with a low activation barrier. Data in c) digitized from reference 73. Knowing these static structural requirements (i.e. wide diffusion pathways, broader bottlenecks for ionic jumps and a large carrier density) has aided in the design and 6 ACS Paragon Plus Environment

Page 7 of 35 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


optimization of many SEs. Starting from α-AgI, Rb+ substitution was used in RbAg4I5 to stabilize the favorable cubic phase at room temperature. Later, the Ag+- and Cu+- conducting cubic Argyrodites were also found. In the LISICON class of Li2+2xZn1-xGeO4, substitution for Ge with Ga, P, As, and V46,47,74–78 helped broaden diffusion pathways, while simultaneously tuning the lithium content. Here, changing the lithium content even alters the conduction mechanism. For instance, the Li-rich compositions exhibit an interstitial mechanism, while Zn-rich Li2+2xZn1-xGeO4 undergoes vacancy-mediated conduction.45,47,78 Still today, the aforementioned SEs are being investigated for the transition between defect-mediated hopping mechanisms and superionic transport, solely due to the diversity of polyanions.79 In a similar fashion, the nature of the polyhedral species has also been shown to affect ionic transport in the lithium thiophosphate glasses.25 The preparation of different solid solutions in NASICON-type SEs (e.g. Li1+xAlxGe2-x(PO4)3 or Na1+xZr2P3-xSixO12), has proven successful in increasing the content of the mobile species as well.32,41,80,81 Changing the lithium concentration in thiophosphate-based classes, like the Argyrodites and Li10GeP2S12, can also tailor the ionic conductivity.24,53,56,82–84 Even the best garnet-structured electrolyte Li7La3Zr2O12 was obtained by increasing the lithium concentration from the initially reported composition of Li5La3M2O12 (M = Nb, Ta).85 In addition to optimizing the carrier concentration, efforts often focus more on broadening the diffusion pathways and opening up the geometric bottlenecks for ion jumps, while ensuring that changes to the local arrangement of the neighboring ions does not disrupt the diffusion pathways.73,86,87 This approach has been successful in the garnet-type ionic conductors,88 the NASICON44,80,89 and LISICON classes,46,47,74–78 as well as lithium and sodium ion conducting thiophosphates.26,27,90,91 While being extensively employed throughout the literature, the general theme of “the broader the diffusion pathways, the lower the activation barriers” emerged.32 However, some materials do not follow this trend. In Li10GeP2S12, isoelectronic substitutions with Si or Sn were used toward obtaining higher Li+ conductivities.56,92,93 While ab initio molecular dynamics simulations predicted that the Si compounds should exhibit higher conductivities,56 the chemical intuition of targeting broader diffusion pathways suggested that the Sn compounds would in fact be better. In order to gain a better understanding of the optimum channel size, Kato et al.94 investigated a series of Li10Ge1−xMxP2S12 (M = Si, Sn) solid solutions and confirmed that the composition with the maximum conductivity and the lowest, most favorable activation barrier is in fact close to the composition of Li10GeP2S12. Thus, despite the predictions, elemental substitutions straying 7 ACS Paragon Plus Environment


Page 8 of 35

from the Li10GeP2S12 composition led to lower ionic conductivities in the end. Figure 3a illustrates the lithium distribution in Li10GeP2S12 as obtained from neutron diffraction data, showing the three-dimensional diffusion of Li+ along the tunnels of the z-axis and in the x-yplane.57 In this structure, the smaller Si4+ leads to a lower conductivity, owing to a geometric restriction of the diffusion pathways.94 However, the structural reasons behind the decrease in conductivity and the concurrent increase in the activation barrier upon moving from Ge4+ to Sn4+ cannot be explained by the ionic radii. As shown in Figure 3c, with increasing unit cell volume and c/a ratio, the activation barrier indeed increases and the ionic conductivity decreases.95 At the same time, a decrease in the S3-S3 distance with increasing unit cell volume can also be observed (Figure 3d). This region of the structure then 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 narrower window, destabilizing the transition state and raising the energetic barrier for ionic motion (Figure 3b). It seems that for the composition of Li10GeP2S12, the interplay between possessing large enough polyhedra and sufficiently broad bottlenecks for ionic jumps is optimal for transport. The structureproperty correlations in Li10GeP2S12 demonstrate that while altering the unit cell volume and diffusion pathways is often a good first approach toward optimizing the transport in 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-plane diffusion does not dominate the lithium mobility in this system, given that an increase in the c/a ratio does not result in an enhancement of the conductivity. Instead, the tunnels along the z-axis are the dominant transport pathways and the in-plane diffusion is likely only used to circumvent mobility issues arising from point-defects, as often observed in one-dimensional ionic conductors.96


Page 9 of 35 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


Figure 3: a) Reconstructed lithium negative nuclear density maps of Li10GeP2S12, illustrating the diffusion along the tunnels in the z-direction and the x-y-plane. b) Li-Li jump in the zdirection with the S3-S3 distance as the geometric bottleneck. c) Increasing activation barrier despite broader diffusion pathways in the x-y-plane and larger c/a ratio in Li10Ge1-xSnxP2S12. d) Interdependence of the bottleneck S3-S3 distance on the conductivity and activation barriers. Figure in a) is adapted with permission from reference 57. Copyright 2016 American Chemical Society. Figures in b)–d) are adapted with permission from reference 95. Copyright 2018 American Chemical Society. The influence of static effects, such as charge carrier density and the breadth of the diffusion pathways, have provided the field with powerful tools for tuning the conductivity of ionic conductors. However, some structures do not allow for such modifications or even behave counterintuitively. Therefore, it is important to combine structural investigations with transport measurements, nuclear magnetic resonance, and theoretical calculations within the field of ionic conductors. 9 ACS Paragon Plus Environment


Page 10 of 35

Future directions will focus on gaining a better understanding of local Coulombic interactions and possible inductive effects in ionic conductors. For instance, in the Li10Ge1-xSnxP2S12 solid solutions, the lower electronegativity and larger ionic radius of Sn4+ leads to less electron density in the M4+ – S2– bonds, as compared to Ge4+.95 With increasing Sn content, the higher electron density on the S2– atoms results in stronger Coulombic attractions between Li+ and S2–, thereby augmenting the activation barrier and the pre-factor 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 inductive effects in solids. In addition to these more bonding- and structurebased approaches, a deeper understanding of correlated ionic motion must also be achieved to better probe the motion of ions in solids.98–100 Again, while often challenging, combined theory and experimental investigations will be required to shed light on how one can initiate and tailor correlated motions, i.e. identify the design principles that activate such correlated motion effects, which may be beneficial for ionic transport.

Figure 4: The activation energy EA and the pre-factor 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-Sbonding characteristics. Figure is adapted with permission from reference 95. Copyright 2018 American Chemical Society.

2.2 The dynamic lattice and “pre-factor dilemma” In addition to the static influences from the underlying crystal structure, increasing the 10 ACS Paragon Plus Environment

Page 11 of 35 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


polarizability of the anion framework has also been suggested to lower the activation barriers for ion jumps.73 Considering Figure 2a, a more polarizable sublattice means softer bonding interactions and a lower energetic cost for the displacement of the lattice during a jump. The idea that the “softness” and dynamics of a host lattice have an influence on the ionic motion in solids had initially been explored in the 1970s/80s,101–105 as the ionic jump processes are thermally activated and must therefore be closely related to the phonon spectrum.106 As the phonon spectrum and its collective vibrational motions directly correlate to the elasticity of the lattice, the covalency and polarizability affect the oscillator strength of the lattice and with it, the transition rate of the hopping ion.105,107,108 Shao-Horn et al. recently highlighted the influence of the phonon frequency (and therefore phonon band centers in the phonon dispersion curve) on the activation barrier, showing that a softer lattice and lower phonon frequencies were correlated with a decrease in the associated activation barriers.109 The expected influence of a softer, more polarizable lattice on the ionic transport is depicted in Figure 5a, in which a softer lattice should lower the activation barrier and the eigenfrequency of the oscillations, i.e. the oscillator or attempt frequency, through a broadening of the local jump oscillators. Beyond the strength of the lattice vibrations, concerted polyhedral rotations have also been shown to assist the mobile ion through a so-called, “paddle wheel mechanism”.110,111 The influence of lattice softening was recently corroborated experimentally using solid solutions of Li6PS5X (X = Cl, Br, I) for Li+ ions and Na3PS4-xSex for Na+ ions, in which the anion polarizability was systematically varied.26,90 In the absence of local structural changes, the activation barrier for ionic motion decreases with decreasing Debye frequency of the lattice (() ), which directly relates to the softness of both the phonon spectrum and the

lattice.26 Figure 5b shows the dependence of the activation energy on a softening lattice in

Na3PS4-xSex. In addition to affecting EA, the pre-factor for ionic motion also decreases over orders of magnitude with a softer lattice. Considering Equation 2, the charge carrier density is not affecting the transport, given the isoelectronic nature of the substitutions, and further, changes in the attempt frequency and jump distance are also not large enough to account for the significant decrease of the pre-factor.26,90 This then leaves the entropy of migration as the main culprit for the detrimental influence of a soft lattice on the transport, 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 approximation, be expressed as:112–114 11 ACS Paragon Plus Environment


/ ∏01 2 342 . ∆ = ln , 01 2 5 6, ∏342 .

Page 12 of 35

Eq. (4)

with the normal frequencies 2 … 1 referring to vibrations around the initial state (I) and the saddle point or transition state (S). While the entropy of migration is affected by nearly all vibrational frequencies around the diffusional pathway, the migration enthalpy is related to the one vibrational mode that carries the ion across the saddle point.47 This attempt frequency is typically approximated by the Debye frequency.26 Therefore, a softening of the lattice leads to lower Debye frequencies (Figure 5b) and is expected to not only affect the activation barrier, but also the pre-factor through a change in both the oscillator frequency and the entropy of migration.26,90,115


Page 13 of 35 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


Figure 5: a) Schematic of the effect of lattice softening on the ionic transport. With increasing softness of the lattice, the local vibrational frequencies decrease along with the activation barrier for the jump. b) Decreasing activation barrier EA and pre-factor with decreasing

Debye frequency 9 in Na3PS4-xSex. c) Meyer-Neldel plot showing the interdependence

between the pre-factor and activation barrier. Figure in a) is adapted with permission from 13 ACS Paragon Plus Environment


reference

26

references

26,90,95

Page 14 of 35

. Copyright 2017 American Chemical Society. Data in b) and c) taken from .


Page 15 of 35 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


Interestingly, the decreasing pre-factor with increasing lattice softness leads to a dilemma, as the approach of introducing more polarizable anions has always been thought to be beneficial for ionic transport. Hence, the influence of the phonon frequencies on the entropy of migration has been largely overlooked and with a decreasing activation barrier, a decrease in the pre-factor is also to be expected (Figure 5c). This is the so-called Meyer-Neldel rule,116– 120

in which the slope of the Meyer-Neldel plot has been linked to the entropy of migration

and the Debye frequencies of the lattice.121 A stiffer lattice may then result in a flatter slope in the Meyer-Neldel plot, while a softer lattice would exhibit a steeper slope. Indeed, for the stiffer Li10Ge1-xSnxP2S12 compounds,95 the slope seems to be flatter relative to the softer Li6PS5X (X = Cl, Br) and Na3PS4-xSex compounds. Thus, materials possessing a flat slope in the Meyer-Neldel plot (i.e. stiffer) are not expected to negatively impact the pre-factor during the optimization of the activation barrier. A notion that is counterintuitive to the known paradigm of “the softer the lattice, the better”. It may then be beneficial to start with stiffer structures and subsequently soften the lattice, instead of starting with soft materials directly, when optimizing ionic transport. On the other hand, in soft materials with a steep slope, chemical modifications toward increasing the pre-factor, such as increased charge carrier densities, may then be less detrimental to the activation barrier. Still, when comparing classes of materials, the more polarizable anion lattices are highly favorable. Moreover, the Debye frequencies of SEs vary with composition26,90,95 and the Meyer-Neldel slope may be decreasing with increasing Debye frequency (i.e. at higher activation barriers). All of the abovementioned considerations will likely be helpful for the optimization of transport in SEs. Accordingly, an alternative approach could be to find structures with soft vibrational modes in the excited state and stiffer phonon modes in the initial ground state, in order to simultaneously obtain high pre-factors and low enthalpies of migration. Both examples, involving the Li6PS5X and Na3PS4-xSex systems, show that the ambiguous influence of lattice dynamics and the “pre-factor dilemma” have been overlooked in the past and that a better understanding of these concepts in relation to ionic transport is required. Further still, recent studies on the influence of a dynamic lattice on the ionic transport have already shown that the entropy of migration and the pre-factor cannot be ignored.122 In particular, when using theoretical approaches to compute ionic conductivities, as well as the data mining for possible structures with high ionic conductivities, it is necessary to include entropic considerations in addition to the calculations of activation barriers.29



2.3 Average and local structures While typical approaches toward enhancing ionic transport involve tailoring the crystal structure through compositional modifications, it is not always clear if the materials are also being influenced by secondary amorphous phases. In particular, in the class of thiophosphate ionic conductors, the glasses and glass ceramics are receiving a lot of attention for this reason.66,67,123–127 The point is that, investigations into material behavior using only Bragg diffraction is proving to be insufficient. Quite often, Raman spectroscopy and nuclear magnetic resonance also need to be used to resolve the different species contributing to the material properties.115,128,129 Dietrich et al. even used pair distribution function (PDF) analysis to show that lithium thiophosphates, which appear crystalline by Bragg diffraction, may actually be “glass-ceramics”.38 These secondary phases, despite revealing no coherent scattering domains,130 may not be entirely amorphous and could be significantly influencing the ionic transport. In poorly conducting crystalline materials, such “glassy” phases may be acting as conducting fillers,38 whereas in good conducting materials, these phases may actually hinder the ionic transport.131 The synthetic conditions must therefore be optimized in order to attain the appropriate balance of glassy and crystalline 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 commonly used. For example, Li10GeP2S12 can only be achieved through such milling techniques. On the other hand, Na3PS4 crystallizes in a tetragonal structure when prepared via classical high-temperature routes,132 but mechanical alloying leads to the cubic polymorph.11,133 Here, the cubic polymorph exhibits orders of magnitude better ionic conductivities than the tetragonal polymorph, despite theoretical predictions suggesting that both should possess similar conductivities.134–136 Using PDF analysis, Krauskopf et al. were able to show that the local structures of both the tetragonal and cubic polymorphs are actually tetragonal (Figure 6). In other words, both polymorphs show the same tetragonal structural motif on the local scale, even though the average crystal structure suggests differences.28 While the typically performed Bragg diffraction only gives a globally averaged depiction of the structure, analysis of the pair distribution function provides a local structural picture, as it represents a histogram of atom-atom distances within the solid in real space. Notably, upon rapidly annealing the high-performance “cubic” polymorph, a tetragonal structure can be obtained, as indicated by Bragg diffraction, however, the transport properties remain 16 ACS Paragon Plus Environment

Page 16 of 35

Page 17 of 35 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


unchanged.28 It is already known that in materials possessing, local dipoles (e.g. ferroelectrics), there can be structural differences between the local and average length scales137–139 and it appears that this is also possible for ionic conductors as well.

Figure 6: Structural representations of the cubic (a) and tetragonal (b) crystal structures, showing the differences in local Na+ ordering. Fitting of the experimentally obtained G(r) data for the cubic polymorph (“c”-Na3PS4) using both the cubic (c) and tetragonal (d) structural models, thereby corroborating 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. 17 ACS Paragon Plus Environment


Page 18 of 35

The differences in the local structure, the enhanced transport arising from processing techniques and the influence of underlying “glassy phases” demonstrate that the structures, compositions and properties of the thiophosphate SEs are all interrelated in a complex manner. In the superionic thiophosphates, the transport behavior does not necessarily depend on the crystal structure alone, but rather the microstructure, the multifaceted nature of the “glass-ceramic” or even defects induced by harsh synthetic conditions. Ultimately, more work must be done in order to answer the many open questions concerning the structure-property relationships at play in these complicated systems.

3. Interfacial processes in solid-state batteries employing thiophosphates The importance of deciphering the principles that govern ionic conduction in solids is clear, given that achieving higher energy density batteries will only be possible through the use of thick electrode configurations coupled with thin separator layers possessing exceptional ionic conductivities.17 Importantly, such thick electrode configurations in SSBs may necessitate even higher ionic conductivities in order to prevent additional cell overpotentials.17 Nevertheless, the ionic conductivity of the SE is not the only bottleneck of the performance for SSBs. During battery operation, reactions between the electrode materials and the electrolyte evolve resistive interfacial layers that strongly affect the subsequent capacity and cycle life. Thus, it is imperative that we dig deeper into both the identity of the interfacial species, as well as their influence on the device properties. 3.1. Interfacial (in)stabilities In viewing the schematic of an archetypical SSB (Figure 1), one can already glean the significant role that interfaces play in influencing the battery performance.140 Importantly, degradation mechanisms of solid electrolytes have been reported throughout the literature regarding both electrode interfaces. On the anode side, the lithium metal electrode reduces many of the best-performing SEs, producing resistive interlayers.3,4,12–15 For example, Wenzel et al. monitored the decomposition of LGPS in contact with Li metal using in situ X-ray photoemission spectroscopy (XPS) and electrochemical impedance spectroscopy. Therein, the continuous growth of an interphase composed of Li2S, reduced Ge species and Li3P was observed, with a corresponding increase in the interfacial resistance.14 While the reaction of 18 ACS Paragon Plus Environment

Page 19 of 35 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


Li metal with various metal-containing solid electrolytes leads to thermodynamically unstable, mixed-conducting interphases that facilitate continuous SE degradation, compositional tailoring of SEs may enable the realization of more favorable, kinetically stable interphases. In other words, one potential approach for mitigating SE degradation could involve the identification of SE compositions that only partially react with Li metal to form thin, ionically conducting, yet electronically insulating interphases.14,141 Additional concerns regarding the anode include dendrite and crack formation, thus making it necessary to find ways to optimize critical current densities.142 Thin film coatings on lithium metal are also being developed to hinder degradation and improve the wetting behavior at the anode interface.143 Meanwhile, on the cathode side, oxidative degradation of thiophosphate SEs against highvoltage active materials is expected4,13 and has also been experimentally corroborated.19,144–151 These reports show that the oxidative SE degradation during cycling stems from the intrinsically narrow electrochemical stability window of the electrolytes. A hypothesis that was further corroborated by first principle calculations from Mo and co-workers.3,4,152 Interfacial reactions can also result in transition metal migration from the cathode active material, hindering the battery performance even more.145 In the end, interfacial reactions involving thiophosphates generate ionically insulating products like Li4P2S7, Li4P2S6 and Li2P2S6.153,154 Therefore, the development of protective coatings for cathode active materials is of paramount importance. Zhang et al. recently extended both the rate capability and the capacity retention in cells constructed with the cathode active material LiCoO2 by employing a LiNb0.5Ta0.5O3 coating.16 The presence of the amorphous oxide layer lowered the interfacial resistance, while also likely inhibiting deleterious interfacial reactions by mitigating direct contact between the cathode material and the SE. Interestingly, these interfacial reactions may not necessarily be irreversible. Redox active behavior of lithium thiophosphates has already been reported in the literature,151,155–158 though an investigation into the active contribution of the SE to the cell performance was still missing. In order to better understand the instability of thiophosphate SEs against high-voltage cathode materials, Koerver et al. used electrochemical impedance spectroscopy and in situ Xray photoemission spectroscopy to demonstrate the redox-activity of β-Li3PS4 induced upon battery operation.18 Notably, these reactions are strongly dependent on the state of charge of the battery, as well as the applied upper cut-off potential. During the charging of the cell, the oxidation of the SE in contact with LiNi0.8Co0.1Mn0.1O2 (NCM-811) can be seen through the 19 ACS Paragon Plus Environment


Page 20 of 35

development of an additional interfacial resistance, relative to the impedance in the discharged state (Figure 7a). With the help of in situ XPS, the evolution of the chemical species could be resolved, thereby explaining the interfacial phenomena (Figure 7b-d). During the degradation of β-Li3PS4, the oxidized products polymerize into interconnected P–[S]n–P type networks, which may lead to the gradual formation of uniquely conducting phases.147,151,156,159 Ultimately, SE degradation at the cathode results in decreased capacity retention through the development of a semi-irreversible cathode/SE interfacial charge transfer resistance, further highlighting the need for suitable protective coatings for the cathode active material. All of the abovementioned conditions occurring at the anode and cathode drive the search for stable (or self-limiting) SE interphases to enable long-term stability and the use of lithium metal anodes.

Figure 7: a) Evolution of the interfacial resistance between the cathode and lithium thiophosphate SE. The resistance is heavily dependent on the state of charge of the battery due to the redox active behavior that can be monitored using in situ X-ray photoemission spectroscopy, shown in b)–d). Figures reproduced with permission from reference

18

.

Copyright 2017 Royal Society of Chemistry.

3.2 Microstructural and volumetric influences In order to electrochemically address all of the active particles simultaneously and avoid local overcharging, achieving adequate percolation and contact within the electrodes is crucial. Solid-state batteries present a different situation from liquid electrolyte-based lithium-ion batteries, where pore diffusion of the liquid electrolyte provides sufficient percolation.5 Thus, added considerations (e.g. particle size/morphology of the active material and composite 20 ACS Paragon Plus Environment

Page 21 of 35 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


electrode composition) must be taken into account when optimizing SSB configurations.16,160 Accordingly, Janek and co-workers were able to show that the ratio of SE to cathode active material in the cathode composite is indeed influential.16 While a high loading of active material leads to higher capacities, the concurrently lower volume fraction of the SE decreases the attainable current densities. It seems to be necessary to design the ratios in the electrode based on the desired application, while also keeping the composite tortuosity in mind.17 For example, thicker electrodes with a higher ratio of cathode active material are required for high-energy applications, while obtaining higher power may require faster ionic conductors or a larger amount of the SE.16,17 Furthermore, good electrode percolation has also been achieved through slurry-based electrode preparations.161–163 More commonly, maintaining sufficient electrode contact in SSBs using thiophosphate-based electrolytes is accomplished through the application of external pressures. Beyond the initial contact, the mechanically soft thiophosphate electrolytes (Young’s modulus of approximately 20 GPa22,164,165) will also undergo microstructural deformations during the cycling-induced expansion and contraction of cathode active materials (e.g. LiCoO2 and LiNi1-x-yCoxMnyO2). Moreover, it was recently shown that despite the soft nature of the thiophosphates, which may allow for the correction of elastic mismatches during cycling, a low fracture toughness exists that may generate additional microstructural flaws.10 Meanwhile, in cathode active materials, the volume changes are structurally related to the deintercalation, the resultant Coulombic repulsion between the MO2 layers and the changing ionic radii of the transition metals.166,167 The crystallographic volume changes occurring within LiCoO2 and NCM-811 during deintercalation are provided in Figure 8a. During charging (i.e. delithiation), the overall volume of the LiCoO2 unit cell expands, while the unit cell volume of NCM-811 contracts. It should also be mentioned that, similar considerations must also be made for the anode side, as most anode materials will experience volume expansion during lithium insertion.20,168 While liquid electrolytes can compensate these volume changes, a tremendous microstructural strain will build up along the grain contacts in SEs.20 Figure 8b shows the nominal strain response upon cycling in an SSB using a metal anode and LiCoO2 as the cathode active material. During charging, the volumes of both the anode and the cathode expand, leading to a reversible pressure increase.20 The resulting pressure change of approximately 1 MPa was found to promote cracking and bending within the SSB due to local microstructural strain.20



Page 22 of 35

When considering volumetric changes, it becomes clear that the pressure inside the cell will vary accordingly. As previously mentioned, for a material like NCM-811, the composite cathode exhibits an overall volume contraction upon delithiation. However, given the reversibility of these volume and pressure changes, one must also consider the microstructural implications within solid-state cells. Figure 8c shows a colorized scanning electron micrograph of NCM-811 after cycling in a solid-state battery. Due to the volume contraction during charge, the spherical particles have lost the intimate contact that they initially had with the SE.19 This contact loss leads to an additional interfacial 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-811 in solid-state batteries,169,170 as compared to liquid electrolyte-based cells (Figure 7d).19 The influence of electrochemically induced local strain must be considered when designing SSBs, especially if they are to be run without external pressure and as such, a certain flexibility needs to be provided,166,167 in order to mitigate volume changes during cycling. It has already been computed that recovering capacity solely by increasing external pressure, thereby reversing contact loss, would not be feasible.171 Therefore, strategies for reducing negative volumetric effects, without employing added external pressure, will be extremely important moving forward.


Page 23 of 35 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


Figure 8: a) Change in unit cell volume during charge (delithiation) of LiCoO2 and NCM811. b) Nominal cell pressure during cycling of an SSB using an In anode and LiCoO2 cathode composite. c) False color scanning electron micrograph showing the contact loss between NCM-811 and the 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 cell using a liquid electrolyte, shown in d). The purple data corresponds to the SSB cell and the orange to the liquid electrolyte cell. Data for Figure a) extrapolated from references

166,167

. b) is

20

reproduced with permission from reference . Copyright 2017 Royal Society of Chemistry. c) and d) are reproduced with permission from reference 19. Copyright 2017 American Chemical Society.

4. Summary and Future Directions In this perspective, we have discussed the structural influences on the ionic conduction in solids, as well as the interfacial chemistry in solid-state batteries. We have highlighted the notion that the typical chemical intuition of attaining broader diffusion pathways may not always be correct and that more in-depth structural investigations are necessary to identify potential bottlenecks for ionic diffusion. We have shown that there is a “pre-factor dilemma” when trying to utilize softer, more polarizable lattices, given the interdependence between the activation barrier and the pre-factor. These results show that the dynamics of the lattice are an important factor governing ionic conduction and that more research is necessary to better understand the associated influences and to find better descriptors for ionic motion in solids. Additionally, we have shown that compositional modifications can alter the electrostatic interactions within the lattice causing inductive effects, which also need to be considered and explored in ionic conductors. We further discussed how the local structures may also differ from the average structures in solid electrolytes, depending on the synthetic conditions, and how this correlates to the transport behavior. Regarding the interfacial processes occurring in SSBs, we have shown that the interfaces and interphases between the electrodes and the SEs must be designed to provide long-term performance and stability. Whether the approach involves protective layers at the anode, and/or thin coatings of high-voltage cathode materials, decomposition of the SE separators needs to be mitigated, if not entirely inhibited. Moreover, we demonstrated that the volumetric changes taking place within the electrodes lead to severe and often detrimental 23 ACS Paragon Plus Environment


Page 24 of 35

microstructural effects that further limit SSB performance. Thus, the optimization of SSB architectures will require the inclusion of both interfacial and microstructural considerations to advance this technology. We hope that this perspective enables future strategies toward understanding ionic conduction in solids and the optimization of solid-state batteries. In the end, the underlying principles governing ionic motion, in addition to the interfacial reactions occurring between the electrodes and the solid electrolyte, must be clarified in order to evaluate and improve soldstate battery technologies in a more realistic fashion.

AUTHOR INFORMATION Corresponding Authors *[email protected]; Notes The authors declare no competing financial interests.

Biographies Sean Culver received his Ph.D. in Inorganic Chemistry in 2016 from the University of Southern California in Los Angeles under the supervision of Prof. Richard Brutchey. He is currently an Alexander von Humboldt postdoctoral fellow at the Justus-Liebig-University Giessen, working with Prof. Jürgen Janek and Dr. Wolfgang Zeier. His research interests include fundamental structure-property relationships in ionic conductors, in addition to the interfacial chemistry and optimization 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 visiting research fellow at the Monash University, Melbourne. Currently he is a Ph.D. candidate at Justus-Liebig-University Giessen under supervision of Prof. Jürgen Janek and Dr. Wolfgang Zeier. His research interests include interfacial reactions and (chemo-)mechanical effects in 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. Jürgen Janek and Dr. Wolfgang Zeier. His research interest includes the


Page 25 of 35 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


fundamental structure-property relationships in ionic conductors and the transfer kinetics between metal anodes and alkali solid electrolytes. Wolfgang Zeier received his Ph.D. in Inorganic Chemistry in 2013 from the JohannesGutenberg University in Mainz under the supervision of Prof. Wolfgang Tremel and Prof. Jeffrey Snyder (California Institute of Technology). After postdoctoral stays at the University of Southern California and at the California Institute of Technology he was appointed group leader at the Justus-Liebig-University Giessen, within the framework of an Emmy-Noether research group. His research interests encompass the fundamental structure-property relationships in solids, with a focus on thermoelectric and ion-conducting materials, as well as solid-solid interfacial chemistry in all-solid-state batteries. Acknowledgements The research was supported by the Deutsche Forschungsgemeinschaft (DFG) under grant number ZE 1010/4-1. S.C. gratefully acknowledges the Alexander von Humboldt Foundation for financial support through a Postdoctoral Fellowship. R.K. gratefully acknowledges financial support by the Funds of the Chemical Industry (FCI).

References (1) (2) (3)

(4)

(5) (6)

(7)

(8)

Goodenough, J. B.; Park, K. S. The Li-Ion Rechargeable Battery: A Perspective. J. Am. Chem. Soc. 2013, 135, 1167–1176. Goodenough, J. B. Rechargeable Batteries: Challenges Old and New. J. Solid State Electrochem. 2012, 16, 2019–2029. Zhu, Y.; He, X.; Mo, Y. Origin of Outstanding Stability in the Lithium Solid Electrolyte Materials: Insights from Thermodynamic Analyses Based on First Principles Calculations. ACS Appl. Mater. Interfaces 2015, 7, 23685–23693. Zhu, Y.; He, X.; Mo, Y. First Principles Study on Electrochemical and Chemical Stability of the Solid Electrolyte-Electrode Interfaces in All-Solid-State Li-Ion Batteries. J. Mater. Chem. A 2015, 4, 3253–3266. Janek, J.; Zeier, W. G. A Solid Future for Battery Development. Nat. Energy 2016, 1, 16141. Kamaya, N.; Homma, K.; Yamakawa, Y.; Hirayama, M.; Kanno, R.; Yonemura, M.; Kamiyama, T.; Kato, Y.; Hama, S.; Kawamoto, K.; Mitsui, A. A Lithium Superionic Conductor. Nat. Mater. 2011, 10, 682–686. Kato, Y.; Hori, S.; Saito, T.; Suzuki, K.; Hirayama, M.; Mitsui, A.; Yonemura, M.; Iba, H.; Kanno, R. High-Power All-Solid-State Batteries Using Sulfide Superionic Conductors. Nat. Energy 2016, 1, 16030. Li, J.; Downie, L. E.; Ma, L.; Qiu, W.; Dahn, J. R. Study of the Failure Mechanisms of 25 ACS Paragon Plus Environment


(9)

(10)

(11)

(12)

(13) (14)

(15)

(16)

(17) (18)

(19)

(20)

(21)

(22) (23)

(24)

LiNi0.8Mn0.1Co0.1O2 Cathode Material for Lithium Ion Batteries. J. Electrochem. Soc. 2015, 162, A1401–A1408. Bucci, G.; Swamy, T.; Chiang, Y.-M.; Carter, W. C. Modeling of Internal Mechanical Failure of All-Solid-State Batteries during Electrochemical Cycling, and Implications for Battery Design. J. Mater. Chem. A 2017, 5, 19422–19430. McGrogan, F. P.; Swamy, T.; Bishop, S. R.; Eggleton, E.; Porz, L.; Chen, X.; Chiang, Y.-M.; Van Vliet, K. J. Compliant Yet Brittle Mechanical Behavior of Li2S-P2S5 Lithium-Ion-Conducting Solid Electrolyte. Adv. Energy Mater. 2017, 7, 1602011. Hayashi, A.; Noi, K.; Sakuda, A.; Tatsumisago, M. Superionic Glass-Ceramic Electrolytes for Room-Temperature Rechargeable Sodium Batteries. Nat. Commun. 2012, 3, 856. Wenzel, S.; Sedlmaier, S. J.; Dietrich, C.; Zeier, W. G.; Janek, J. Interfacial Reactivity and Interphase Growth of Argyrodite Solid Electrolytes at Lithium Metal Electrodes. Solid State Ionics 2018, 318, 102–112. Richards, W. D.; Miara, L. J.; Wang, Y.; Kim, J. C.; Ceder, G. Interface Stability in Solid-State Batteries. Chem. Mater. 2016, 28, 266–273. Wenzel, S.; Randau, S.; Leichtweiß, T.; Weber, D. A.; Sann, J.; Zeier, W. G.; Janek, J. Direct Observation of the Interfacial Instability of the Fast Ionic Conductor Li10GeP2S12 at the Lithium Metal Anode. Chem. Mater. 2016, 28, 2400–2407. Wenzel, S.; Leichtweiss, T.; Weber, D. A.; Sann, J.; Zeier, W. G.; Janek, J. Interfacial Reactivity Benchmarking of the Sodium Ion Conductors Na3PS4 and Sodium βAlumina for Protected Sodium Metal Anodes and Sodium All-Solid-State Batteries. ACS Appl. Mater. Interfaces 2016, 8, 28216–28224. Zhang, W.; Weber, D. A.; Weigand, H.; Arlt, T.; Manke, I.; Schröder, D.; Koerver, R.; Leichtweiss, T.; Hartmann, P.; Zeier, W. G.; Janek, J. Interfacial Processes and Influence of Composite Cathode Microstructure Controlling the Performance of AllSolid-State Lithium Batteries. ACS Appl. Mater. Interfaces 2017, 9, 17835–17845. Kato, Y.; Shiotani, S.; Morita, K.; Suzuki, K.; Hirayama, M.; Kanno, R. All-Solid-State Batteries with Thick Electrode Configurations. J. Phys. Chem. Lett. 2018, 607–613. Koerver, R.; Walther, F.; Aygün, I.; Sann, J.; Dietrich, C.; Zeier, W. G.; Janek, J. Redox-Active Cathode Interphases in Solid-State Batteries. J. Mater. Chem. A 2017, No. 43, 22750–22760. Koerver, R.; Aygün, I.; Leichtweiß, T.; Dietrich, C.; Zhang, W.; Binder, J. O.; Hartmann, P.; Zeier, W. G.; Janek, J. Capacity Fade in Solid-State Batteries: Interphase Formation and Chemomechanical Processes in Nickel-Rich Layered Oxide Cathodes and Lithium Thiophosphate Solid Electrolytes. Chem. Mater. 2017, 29, 5574–5582. Zhang, W.; Schröder, D.; Arlt, T.; Manke, I.; Koerver, R.; Pinedo, R.; Weber, D. A.; Sann, J.; Zeier, W. G.; Janek, J. (Electro)chemical Expansion during Cycling: Monitoring the Pressure Changes in Operating Solid-State Lithium Batteries. J. Mater. Chem. A 2017, 5, 9929–9936. Deng, Z.; Wang, Z.; Chu, I.-H.; Luo, J.; Ong, S. P. Elastic Properties of Alkali Superionic Conductor Electrolytes from First Principles Calculations. J. Electrochem. Soc. 2016, 163, A67–A74. Sakuda, A.; Hayashi, A.; Tatsumisago, M. Sulfide Solid Electrolyte with Favorable Mechanical Property for All-Solid-State Lithium Battery. Sci. Rep. 2013, 3, 2261. Sakuda, A.; Hayashi, A.; Takigawa, Y.; Higashi, K.; Tatsumisago, M. Evaluation of Elastic Modulus of Li2S–P2S5 Glassy Solid Electrolyte by Ultrasonic Sound Velocity Measurement and Compression Test. J. Ceram. Soc. Jap. 2013, 121, 946–949. Kuhn, A.; Gerbig, O.; Zhu, C.; Falkenberg, F.; Maier, J.; Lotsch, B. V. A New Ultrafast Superionic Li-Conductor: Ion Dynamics in Li11Si2PS12 and Comparison with 26 ACS Paragon Plus Environment

Page 26 of 35

Page 27 of 35 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


(25)

(26)

(27)

(28) (29)

(30) (31) (32) (33) (34) (35) (36) (37) (38)

(39) (40) (41)

(42) (43)

(44)

Other Tetragonal LGPS-Type Electrolytes. Phys. Chem. Chem. Phys. 2014, 16, 14669– 14674. Dietrich, C.; Weber, D.; Sedlmaier, S. J.; Indris, S.; Culver, S.; Walter, D.; Janek, J.; Zeier, W. Lithium Ion Conductivity in Li2S-P2S5 Glasses – Building Units and Local Structure Evolution during the Crystallization of the Superionic Conductors Li3PS4 , Li7P3S11 . J. Mater. Chem. A 2017, 5, 18111–18119. Kraft, M. A.; Culver, S. P.; Calderon, M.; Böcher, F.; Krauskopf, T.; Senyshyn, A.; Dietrich, C.; Zevalkink, A.; Janek, J.; Zeier, W. G. Influence of Lattice Polarizability on the Ionic Conductivity in the Lithium Superionic Argyrodites Li6PS5X (X = Cl, Br, I). J. Am. Chem. Soc. 2017, 139, 10909–10918. Minafra, N.; Culver, S. P.; Krauskopf, T.; Senyshyn, A.; Zeier, W. G. Effect of Si Substitution on the Structural and Transport Properties of Superionic Li-Argyrodites. J. Mater. Chem. A 2018, 6, 645–651. Krauskopf, T.; Culver, S. P.; Zeier, W. G. Local Tetragonal Structure of the Cubic Superionic Conductor Na33PS4. Inorg. Chem. 2018, 57, 4739–4744. Sendek, A. D.; Yang, Q.; Cubuk, E. D.; Duerloo, K.-A. N.; Cui, Y.; Reed, E. J. Holistic Computational Structure Screening of More than 12 000 Candidates for Solid LithiumIon Conductor Materials. Energy Environ. Sci. 2017, 10, 306–320. Sato, H. Some Theoretical Aspects of Solid Electrolytes. In Solid Electrolytes; Geller, S., Ed.; Springer Verlag: Berlin, 1977; pp 3–39. Boyce, J. B.; Huberman, B. A. Superionic Conductors: Transitions, Structures, Dynamics. Phys. Rep. 1979, 51, 189–265. West, A. R. Basic Solid State Chemistry, 2nd Edition; John Wiley & Sons: West Sussex, England, 2008. Tarascon, J. M.; Armand, M. Issues and Challenges Facing Rechargeable Lithium Batteries. Nature 2001, 414, 359–367. Goodenough, J. Ceramic Solid Electrolytes. Solid State Ionics 1997, 94, 17–25. Funke, K. Jump Relaxation in Solid Ionic Conductors. Solid State Ionics 1993, 22, 111–195. Funke, K. Debye-Hückel-Type Relaxation Processes in Solid Ionic Conductors: The Model. Solid State Ionics 1986, 18, 183–190. Tilley, R. J. D. Defects in Solids, 10th ed.; Wiley VCH: Hoboken, New Jersey, 2008. Dietrich, C.; Sadowski, M.; Sabrina, S.; Weber, D. A.; Sedlmaier, S. J.; Weldert, K. S.; Indris, S.; Albe, K.; Janek, J.; Zeier, W. G. Local Structural Investigations, Defect Formation and Ionic Conductivity of the Lithium Ionic Conductor Li4P2S6. Chem. Mater. 2016, 28, 8764–8773. Janek, J.; Martin, M. Electrotransport in Ionic Crystals: II. A Dynamical Model. Ber. Bunsenges. Phys. Chemie 1994, 98, 665–673. Guin, M.; Tietz, F. Survey of the Transport Properties of Sodium Superionic Conductor Materials for Use in Sodium Batteries. J. Pow. Source 2015, 273, 1056–1064. Francisco, B. E.; Stoldt, C. R.; M’Peko, J. C. Lithium-Ion Trapping from Local Structural Distortions in Sodium Super Ionic Conductor (NASICON) Electrolytes. Chem. Mater. 2014, 26, 4741–4749. Francisco, B. E.; Stoldt, C. R.; M’Peko, J.-C. Energetics of Ion Transport in NASICON-Type Electrolytes. J. Phys. Chem. C 2015, 119, 16432–16442. Monchak, M.; Hupfer, T.; Senyshyn, A.; Boysen, H.; Chernyshov, D.; Hansen, T.; Schell, K. G.; Bucharsky, E. C.; Hoffmann, M. J.; Ehrenberg, H. Lithium Diffusion Pathway in Li1.3Al0.3Ti1.7(PO4)3(LATP) Superionic Conductor. Inorg. Chem. 2016, 55, 2941–2945. Losilla, E. R.; Aranda, M. A. G.; Bruque, S.; París, M. A.; Sanz, J.; West, A. R. 27 ACS Paragon Plus Environment


(45)

(46)

(47) (48) (49) (50) (51) (52) (53)

(54)

(55) (56)

(57)

(58)

(59)

(60)

(61)

Understanding Na Mobility in NASICON Materials: A Rietveld, 23Na and 31P MAS NMR, and Impedance Study. Chem. Mater. 1998, 10, 665–673. Abrahams, I.; Bruce, P. G.; West, A. R.; David, W. I. F. Structure Determination of LISICON Solid Solutions by Powder Neutron Diffraction. J. Solid State Chem. 1988, 75, 390–396. Robertson, A.; West, A. R. Phase Equilibria, Crystal Chemistry and Ionic Conductivity in the LISICON System Li4GeO4Li2.5Ga0.5GeO4. Solid State Ionics 1992, 58 (3–4), 351–358. Bruce, P. G.; West, A. R. Ionic Conductivity of LISICON Solid Solutions, Li2+2xZn1xGeO4. J. Solid State Chem. 1982, 44, 354–365. Thangadurai, V.; Weppner, W. Recent Progress in Solid Oxide and Lithium Ion Conducting Electrolytes Research. Ionics 2006, 12, 81–92. Thangadurai, V.; Kaack, H.; Weppner, W. Novel Fast Lithium Ion Conduction in Garnet-Type Li5La3M2O12 (M = Nb, Ta). J. Am. Ceram. Soc. 2003, 86, 437. Murugan, R.; Thangadurai, V.; Weppner, W. Fast Lithium Ion Conduction in GarnetType Li7La3Zr2O12. Angew. Chemie - Int. Ed. 2007, 46, 7778–7781. Thangadurai, V.; Weppner, W. Li6ALa2Ta2O12 (A = Sr, Ba): Novel Garnet-Like Oxides for Fast Lithium Ion Conduction. Adv. Funct. Mater. 2005, 15, 107–112. Adams, S.; Prasada Rao, R. Structural Requirements for Fast Lithium Ion Migration in Li10GeP2S12. J. Mater. Chem. 2012, 22, 7687–7691. Hori, S.; Suzuki, K.; Hirayama, M.; Kato, Y.; Saito, T.; Yonemura, M.; Kanno, R. Synthesis, Structure, and Ionic Conductivity of Solid Solution, Li10+δM1+δP2−δS12 (M = Si, Sn). Faraday Discuss. 2014, 176, 83–94. Hori, S.; Taminato, S.; Suzuki, K.; Hirayama, M.; Kato, Y.; Kanno, R. Structure– property Relationships in Lithium Superionic Conductors Having a Li10GeP2S12 -Type Structure. Acta Cryst. B Eng. Mater. 2015, 71, 727–736. Mo, Y.; Ong, S. P.; Ceder, G. First Principles Study of the Li10GeP2S12 Lithium Super Ionic Conductor Material. Chem. Mater. 2012, 24, 15–17. Ong, S. P.; Mo, Y.; Richards, W. D.; Miara, L.; Lee, H. S.; Ceder, G. Phase Stability, Electrochemical Stability and Ionic Conductivity of the Li10±1MP2X12 (M = Ge, Si, Sn, Al or P, and X = O, S or Se) Family of Superionic Conductors. Energy Environ. Sci. 2013, 6, 148–156. Weber, D. A.; Senyshyn, A.; Weldert, K. S.; Wenzel, S.; Zhang, W.; Kaiser, R.; Berendts, S.; Janek, J.; Zeier, W. G. Structural Insights and 3D Diffusion Pathways within the Lithium Superionic Conductor Li10GeP2S12. Chem. Mater. 2016, 28, 5905−5915. Zhu, Z.; Chu, I. H.; Ong, S. P. Li3Y(PS4)2 and Li5PS4Cl2: New Lithium Superionic Conductors Predicted from Silver Thiophosphates Using Efficiently Tiered Ab Initio Molecular Dynamics Simulations. Chem. Mater. 2017, 29, 2474–2484. de Klerk, N. J. J.; Rosłoń, I.; Wagemaker, M. Diffusion Mechanism of Li Argyrodite Solid Electrolytes for Li-Ion Batteries and Prediction of Optimized Halogen Doping: The Effect of Li Vacancies, Halogens, and Halogen Disorder. Chem. Mater. 2016, 28, 7955–7963. Deiseroth, H. J.; Kong, S. T.; Eckert, H.; Vannahme, J.; Reiner, C.; Zaiß, T.; Schlosser, M. Li6PS5X: A Class of Crystalline Li-Rich Solids with an Unusually High Li+ Mobility. Angew. Chem. Int. Ed. 2008, 47, 755–758. Deiseroth, H.; Maier, J.; Weichert, K.; Nickel, V.; Kong, S.; Reiner, C. Li7PS6 and Li6PS5X (X: Cl, Br, I): Possible Three-Dimensional Diffusion Pathways for Lithium Ions and Temperature Dependence of the Ionic Conductivity by Impedance Measurements. Z. Anorg. Allg. Chem. 2011, 637, 1287–1294. 28 ACS Paragon Plus Environment

Page 28 of 35

Page 29 of 35 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


(62)

(63)

(64)

(65)

(66)

(67)

(68) (69) (70)

(71)

(72) (73)

(74) (75)

(76) (77) (78) (79)

(80)

Holzmann, T.; Schoop, L. M.; Ali, M. N.; Moudrakovski, I.; Gregori, G.; Maier, J.; Cava, R. J.; Lotsch, B. V. Li0.6[Li0.2Sn0.8S2] – a Layered Lithium Superionic Conductor. Energy Environ. Sci. 2016, 9, 2578–2585. Duchardt, M.; Ruschewitz, U.; Adams, S.; Dehnen, S.; Roling, B. Vacancy-Controlled Na+ Superion Conduction in Na11Sn2PS12. Angew. Chemie Int. Ed. 2018, 57, 1351– 1355. Richards, W. D.; Tsujimura, T.; Miara, L. J.; Wang, Y.; Kim, J. C.; Ong, S. P.; Uechi, I.; Suzuki, N.; Ceder, G. Design and Synthesis of the Superionic Conductor Na10SnP2S12. Nat. Commun. 2016, 7, 11009. Zhang, Z.; Ramos, E.; Lalère, F.; Assoud, A.; Kaup, K.; Hartman, P.; Nazar, L. F. Na11Sn2PS12 : A New Solid State Sodium Superionic Conductor. Energy Environ. Sci. 2018, 11, 87–93. Busche, M. R.; Weber, D. A.; Schneider, Y.; Dietrich, C.; Wenzel, S.; Leichtweiss, T.; Schröder, D.; Zhang, W.; Weigand, H.; Walter, D.; Sedlmaier, S. J.; Houtarde, D.; Nazar, L. F.; Janek, J. In Situ Monitoring of Fast Li-Ion Conductor Li7P3S11 Crystallization Inside a Hot-Press Setup. Chem. Mater. 2016, 28, 6152–6165. Seino, Y.; Nakagawa, M.; Senga, M.; Higuchi, H.; Takada, K.; Sasaki, T. Analysis of the Structure and Degree of Crystallisation of 70Li2S–30P2S5 Glass Ceramic. J. Mater. Chem. A 2015, 3, 2756–2761. Owens, B. B.; Argue, G. R. High-Conductivity Solid Electrolytes: MAg4I5. Science 1967, 157, 308–210. Geller, S. Crystal Structure of the Solid Electrolyte, RbAg4I5. Science 1967, 157, 310– 312. Nilges, T.; Pfitzner, A. A Structural Differentiation of Quaternary Copper Argyrodites: Structure - Property Relations of High Temperature Ion Conductors. Zeitschrift fur Krist. 2005, 220, 281–294. Gaudin, E.; Petricek, V.; Boucher, F.; Taulelle, F.; Evain, M. Structures and Phase Transitions of the A7PSe6 (A = Ag, Cu) Argyrodite-Type Ionic Conductors. III. αCu7PSe6. Acta Crystallogr. Sect. B Struct. Sci. 2000, 56, 972–979. Rice, M. J.; Roth, W. L. Ionic Transport in Superionic Conductors: A Theoretical Model. J. Solid State Chem. 1972, 4, 294–310. Wang, Y.; Richards, W. D.; Ong, S. P.; Miara, L. J.; Kim, J. C.; Mo, Y.; Ceder, G. Design Principles for Solid-State Lithium Superionic Conductors. Nat. Mater. 2015, 14, 1026–1031. Hu, Y.-W. Ionic Conductivity of Lithium Orthosilicate—Lithium Phosphate Solid Solutions. J. Electrochem. Soc. 1977, 124, 1240. Rodger, A. Li+ Ion Conducting γ Solid Solutions in the Systems Li4XO4-LiYO4: X=Si, Ge, Ti; Y=P, As, V; Li4XO4-LiZO2: Z=Al, Ga, Cr and Li4GeO4-Li2CaGeO4. Solid State Ionics 1985, 15, 185–198. Kuwano, J.; West, A. R. New Li+ Ion Conductors in the System, Li4GeO4-Li3VO4. Mater. Res. Bull. 1980, 15, 1661–1667. Kamphorst, J.; Hellstrom, E. Fast Li Ionic Conduction in Solid Solutions of the System Li4GeO4-Li2ZnGeO4-Li3PO4. Solid State Ionics 1980, 1, 187–197. Bruce, P. G.; West, A. R. Phase Diagram of Lisicon. MRS Bull. 1980, 15, 379–385. Deng, Y.; Eames, C.; Fleutot, B.; David, R.; Chotard, J. N.; Suard, E.; Masquelier, C.; Islam, M. S. Enhancing the Lithium Ion Conductivity in Lithium Superionic Conductor (LISICON) Solid Electrolytes through a Mixed Polyanion Effect. ACS Appl. Mater. Int. 2017, 9, 7050–7058. Guin, M.; Tietz, F. Survey of the Transport Properties of Sodium Superionic Conductor Materials for Use in Sodium Batteries. J. Power Sources 2015, 273, 1056–1064. 29 ACS Paragon Plus Environment


(81)

Weiss, M.; Weber, D. A.; Senyshyn, A.; Janek, J.; Zeier, W. G. Correlating Transport and Structural Properties in Li1+xAlxGe2-x(PO4)3 (LAGP) Prepared from Aqueous Solution. ACS Appl. Mater. Interfaces 2018, 10, 10935–10944. (82) Hood, Z. D.; Kates, C.; Kirkham, M.; Adhikari, S.; Liang, C.; Holzwarth, N. A. W. Structural and Electrolyte Properties of Li4P2S6. Solid State Ionics 2016, 284, 61–70. (83) Kuhn, A.; Duppel, V.; Lotsch, B. V. Tetragonal Li10GeP2S12 and Li7GePS8 – Exploring the Li Ion Dynamics in LGPS Li Electrolytes. Energy Environ. Sci. 2013, 6, 3548– 3552. (84) Kwon, O.; Hirayama, M.; Suzuki, K.; Kato, Y.; Saito, T.; Yonemura, M.; Kamiyama, T.; Kanno, R. Synthesis, Structure, and Conduction Mechanism of the Lithium Superionic Conductor Li10+δGe1+δP2−δS12. J. Mater. Chem. A 2015, 3, 438–446. (85) Zeier, W. G. Structural Limitations for Optimizing Garnet-Type Solid Electrolytes: A Perspective. Dalt. Trans. 2014, 43, 16133–16138. (86) Urban, A.; Lee, J.; Ceder, G. The Configurational Space of Rocksalt-Type Oxides for High-Capacity Lithium Battery Electrodes. Adv. Energy Mater. 2014, 4, 1400478. (87) Mahan, G. D. Theoretical Issues in Superionic Conductors. In Superionic Conductors; Mahan, G. D., Ed.; Springer, 1976; pp 115–134. (88) Zeier, W. G.; Zhou, S.; Lopez-Bermudez, B.; Page, K.; Melot, B. C. Dependence of the Li-Ion Conductivity and Activation Energies on the Crystal Structure and Ionic Radii in Li6MLa2Ta2O12. ACS Appl. Mater. Interface 2014, 6, 10900–10907. (89) Lang, B.; Ziebarth, B.; Elsässer, C. Lithium Ion Conduction in LiTi2(PO4)3 and Related Compounds Based on the NASICON Structure: A First-Principles Study. Chem. Mater. 2015, 27, 5040–5048. (90) Krauskopf, T.; Pompe, C.; Kraft, M. A.; Zeier, W. G. Influence of Lattice Dynamics on Na+Transport in the Solid Electrolyte Na3PS4-xSex. Chem. Mater. 2017, 29, 8859–8869. (91) Zhang, L.; Yang, K.; Mi, J.; Lu, L.; Zhao, L.; Wang, L.; Li, Y.; Zeng, H. Na3PSe4: A Novel Chalcogenide Solid Electrolyte with High Ionic Conductivity. Adv. Energy Mater. 2015, 5, 1501294. (92) Bron, P.; Johansson, S.; Zick, K.; Schmedt auf der Günne, J.; Dehnen, S.; Roling, B. Li10SnP2S12: An Affordable Lithium Superionic Conductor. J. Am. Chem. Soc. 2013, 135, 15694–15697. (93) Wang, H.; Cao, X.; Takagiwa, Y.; Snyder, G. J. Higher Mobility in Bulk Semiconductors by Separating the Dopants from the Charge-Conducting Band – a Case Study of Thermoelectric PbSe. Mater. Horizons 2015, 2, 323–329. (94) Kato, Y.; Saito, R.; Sakano, M.; Mitsui, A.; Hirayama, M.; Kanno, R. Synthesis, Structure and Lithium Ionic Conductivity of Solid Solutions of Li10(Ge1-xMx)P2S12(M = Si, Sn). J. Power Sources 2014, 271, 60–64. (95) Krauskopf, T.; Culver, S. P.; Zeier, W. G. Bottleneck of Diffusion and Inductive Effects in Li10Ge1-xSnxP2S12. Chem. Mater. 2018, 30, 1791–1798. (96) Morgan, D.; Van der Ven, A.; Ceder, G. Li Conductivity in LixMPO4 (M = Mn, Fe, Co, Ni) Olivine Materials. Electrochem. Solid-State Lett. 2004, 7, A30–A32. (97) Yu, Z.; Shang, S.-L.; Seo, J.-H.; Wang, D.; Luo, X.; Huang, Q.; Chen, S.; Lu, J.; Li, X.; Liu, Z.-K.; Wang, D. Exceptionally High Ionic Conductivity in Na3P0.62As0.38S4 with Improved Moisture Stability for Solid-State Sodium-Ion Batteries. Adv. Mater. 2017, 29, 1605561. (98) He, X.; Zhu, Y.; Mo, Y. Conductors. Nat. Commun. 2017, 8, 15893. (99) Morgan, B. J. Lattice-Geometry Effects in Garnet Solid Electrolytes: A Lattice-Gas Monte Carlo Simulation Study. R. Soc. open sci. 2017, 4, 170824. (100) Marcolongo, A.; Marzari, N. Ionic Correlations and Failure of Nernst-Einstein Relation in Solid-State Electrolytes. Phys. Rev. Mater. 2017, 1, 25402. 30 ACS Paragon Plus Environment

Page 30 of 35

Page 31 of 35 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


(101) Sen, P. N.; Huberman, B. A. Low-Frequency Response of Superionic Conductors. Phys. Rev. Lett. 1975, 34, 1059–1060. (102) Huberman, B.; Sen, P. Dielectric Response of a Superionic Conductor. Phys. Rev. Lett. 1974, 33, 1379–1382. (103) Allen, S. J.; Remeika, J. P. Direct Measurement of the Attempt Frequency for Ion Diffusion in Ag and Na β-Alumina. Phys. Rev. Lett. 1974, 33, 1478–1481. (104) Pardee, W. J.; Mahan, G. D.; Eastman, D. E.; Pollak, R. A.; Ley, L.; McFeely, F. R.; Kowalczyk, S. P.; Shirley, D. A. Analysis Of Surface-Plasmon And Bulk-Plasmon Contributions To X-Ray Photoemission Spectra. Phys. Rev. B 1975, 11, 3614–3616. (105) Zeller, H. R.; Brüesch, P.; Pietronero, L.; Strässler, S. Lattice Dynamics and Ionic Motion. In Superionic Conductors; Mahan, G. D., Ed.; Springer, 1976; pp 201–214. (106) Bührer, W.; Nicklow, R. M.; Brüesch, P. Lattice Dynamics of β-(Silver Iodide) by Neutron Scattering. Phys. Rev. B 1978, 17, 3362–3370. (107) Bruesch, P.; Pietronero, L.; Strassler, S.; Zeller, H. R. Brownian Motion in a Polarizable Lattice: Application to Superionic Conductors. Phys. Rev. B 1977, 15, 4631–4637. (108) Fischer, K.; Bilz, H.; Haberkorn, R.; Weber, W. Covalency and Deformability of Ag+‐ Ions in the Lattice Dynamics of Silver Halides. Phys. status solidi 1972, 54, 285–294. (109) Bachman, J. C.; Muy, S.; Grimaud, A.; Chang, H.-H.; Pour, N.; Lux, S. F.; Paschos, O.; Maglia, F.; Lupart, S.; Lamp, P.; Giordano, L.; Shao-Horn, Y. Inorganic SolidState Electrolytes for Lithium Batteries: Mechanisms and Properties Governing Ion Conduction. Chem. Rev. 2016, 116, 140–162. (110) Jansen, M. Volume Effect or Paddle-Wheel Mechanism—Fast Alkali-Metal Ionic Conduction in Solids with Rotationally Disordered Complex Anions. Angew. Chemie Int. Ed. English 1991, 30, 1547–1558. (111) Li, X.; Benedek, N. A. Enhancement of Ionic Transport in Complex Oxides through Soft Lattice Modes and Epitaxial Strain. Chem. Mater. 2015, 27, 2647–2652. (112) Toyoura, K.; Koyama, Y.; Kuwabara, A.; Oba, F.; Tanaka, I. First-Principles Approach to Chemical Diffusion of Lithium Atoms in a Graphite Intercalation Compound. Phys. Rev. B 2008, 79, 214303. (113) Vineyard, G. H. Frequency Factors and Isotope Effects in Solid State Rate Processes. J. Phys. Chem Solids 1957, 3, 121–127. (114) Mehrer, H. Diffusion in Solids: Fundamentals, Methods, Materials, DiffusionControlled Processes; Springer-Verlag Berlin Heidelberg, 2007. (115) Breuer, S.; Wilkening, M. Mismatch in Cation Size Causes Rapid Anion Dynamics in Solid Electrolytes: The Role of the Arrhenius Pre-Factor. Dalt. Trans. 2018, 47, 4105– 4117. (116) Kreuer, K.-D.; Kohler, H.; Maier, J. Sodium Ion Conductors with NASICON Framework Structure. In High Conductivity Ionic Conductors: Recent Trends and Applications; Takahash, T., Ed.; World Scientific, 1989; pp 242–279. (117) Meyer, W.; Neldel, H. Relation between the Energy Constant and the Quantity Constant in the Conductivity–temperature Formula of Oxide Semiconductors. Z. Tech. Phys. 1937, 588. (118) Ngai, K. Meyer-Neldel Rule and Anti Meyer-Neldel Rule of Ionic Conductivity Conclusions from the Coupling Model. Solid State Ionics 1998, 105, 231–235. (119) Nowick, A. S.; Lee, W. K.; Jain, H. Survey and Interpretation of Pre-Exponentials of Conductivity. Solid State Ionics 1988, 28–30, 89–94. (120) Knödler, D.; Pendzig, P.; Dieterich, W. Ion Dynamics in Structurally Disordered Materials: Effects of Random Coulombic Traps. Solid State Ionics 1996, 86–88, 29–39. (121) Yelon, A.; Movaghar, B. Reply to “Comment on ‘Origin and Consequences of the 31 ACS Paragon Plus Environment


Compensation (Meyer-Neldel) Law’ .” Phys. Rev. B 1992, 46, 12244–12250. (122) Muy, S.; Bachman, J. C.; Giordano, L.; Chang, H.-H.; Abernathy, D. L.; Bansal, D.; Delaire, O.; Hori, S.; Kanno, R.; Maglia, F.; Lupart, S.; Lamp, P.; Shao-Horn, Y. Tuning Mobility and Stability of Lithium Ion Conductors Based on Lattice Dynamics. Energy Environ. Sci. 2018, 11, 850–859. (123) Tatsumisago, M.; Hayashi, A. Sulfide Glass-Ceramic Electrolytes for All-Solid-State Lithium and Sodium Batteries. Int. J. Appl. Glas. Sci. 2014, 5, 226–235. (124) Tatsumisago, M.; Mizuno, F.; Hayashi, A. All-Solid-State Lithium Secondary Batteries Using Sulfide-Based Glass–ceramic Electrolytes. J. Power Sources 2006, 159, 193– 199. (125) Seino, Y.; Ota, T.; Takada, K. A Sulphide Lithium Super Ion Conductor Is Superior to Liquid Ion Conductors for Use in Rechargeable Batteries. Energy Environ. Sci. 2014, 7, 627. (126) Murakami, M.; Shimoda, K.; Shiotani, S.; Mitsui, A.; Ohara, K.; Onodera, Y.; Arai, H.; Uchimoto, Y.; Ogumi, Z. Dynamical Origin of Ionic Conductivity for Li7P3S11 Metastable Crystal As Studied by 6,7Li and 31P Solid-State NMR. J. Phys. Chem. C 2015, 119, 24248–24254. (127) Minami, T.; Hayashi, A.; Tatsumisago, M. Recent Progress of Glass and GlassCeramics as Solid Electrolytes for Lithium Secondary Batteries. Solid State Ionics 2006, 177, 2715–2720. (128) Epp, V.; Gün, Ö.; Deiseroth, H. J.; Wilkening, M. Highly Mobile Ions: LowTemperature NMR Directly Probes Extremely Fast Li+ Hopping in Argyrodite-Type Li6PS5Br. J. Phys. Chem. Lett. 2013, 4, 2118–2123. (129) Wohlmuth, D.; Epp, V.; Wilkening, M. Fast Li Ion Dynamics in the Solid Electrolyte Li7P3S11 as Probed by 6,7Li NMR Spin-Lattice Relaxation. ChemPhysChem 2015, 16, 2582–2593. (130) Egami, T.; Billinge, S. J. L. Underneath the Bragg Peaks: Structural Analysis of Complex Materials, 7th ed.; Cahn, R. W., Ed.; Pergamon Materials Series: Oxford, 2003. (131) Tsukasaki, H.; Mori, S.; Shiotani, S.; Yamamura, H.; Iba, H. Direct Observation of a Non-Isothermal Crystallization Process in Precursor Li10GeP2S12 Glass Electrolyte. J. Power Sources 2017, 369, 57–64. (132) Jansen, M.; Henseler, U. Synthesis, Structure Determination, and Ionic Conductivity of Sodium Tetrathiophosphate. J. Solid State Chem. 1992, 99, 110–119. (133) Hayashi, A.; Noi, K.; Tanibata, N.; Nagao, M.; Tatsumisago, M. High Sodium Ion Conductivity of Glass-Ceramic Electrolytes with Cubic Na3PS4. J. Pow. Source 2014, 258, 420–423. (134) Yu, C.; Ganapathy, S.; de Klerk, N. J. J.; van Eck, E. R. H.; Wagemaker, M. Na-Ion Dynamics in Tetragonal and Cubic Na3PS4, a Na-Ion Conductor for Solid State Na-Ion Batteries. J. Mater. Chem. A 2016, 4, 15095–15105. (135) Zhu, Z.; Chu, I. H.; Deng, Z.; Ong, S. P. Role of Na+ Interstitials and Dopants in Enhancing the Na+ Conductivity of the Cubic Na3PS4 Superionic Conductor. Chem. Mater. 2015, 27, 8318–8325. (136) De Klerk, N. J. J.; Wagemaker, M. Diffusion Mechanism of the Sodium-Ion Solid Electrolyte Na3PS4 and Potential Improvements of Halogen Doping. Chem. Mater. 2016, 28, 3122–3130. (137) Aksel, E.; Forrester, J. S.; Nino, J. C.; Page, K.; Shoemaker, D. P.; Jones, J. L. Local Atomic Structure Deviation from Average Structure of Na0.5Bi0.5TiO3: Combined XRay and Neutron Total Scattering Study. Phys. Rev. B 2013, 87, 104113. (138) Page, K.; Proffen, T.; Niederberger, M.; Seshadri, R. Probing Local Dipoles and 32 ACS Paragon Plus Environment

Page 32 of 35

Page 33 of 35 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


Ligand Structure in BaTiO3 Nanoparticles. Chem. Mater. 2010, 22, 4386–4391. (139) Page, K.; Stoltzfus, M. W.; Kim, Y.-I.; Proffen, T.; Woodward, P. M.; Cheetham, A. K.; Seshadri, R. Local Atomic Ordering in BaTaO2N Studied by Neutron Pair Distribution Function Analysis and Density Functional Theory. Chem. Mater. 2007, 19, 4037–4042. (140) Luntz, A. C.; Voss, J.; Reuter, K. Interfacial Challenges in Solid-State Li Ion Batteries. J. Phys. Chem. Lett. 2015, 6, 4599–4604. (141) Kato, A.; Kowada, H.; Deguchi, M.; Hotehama, C.; Hayashi, A.; Tatsumisago, M. XPS and SEM Analysis between Li/Li3PS4 Interface with Au Thin Film for All-Solid-State Lithium Batteries. Solid State Ionics 2018, 322, 1–4. (142) Porz, L.; Swamy, T.; Sheldon, B. W.; Rettenwander, D.; Frömling, T.; Thaman, H. L.; Berendts, S.; Uecker, R.; Carter, W. C.; Chiang, Y.-M. M. Mechanism of Lithium Metal Penetration through Inorganic Solid Electrolytes. Adv. Energy Mater. 2017, 7, 1– 12. (143) Zhang, Z.; Chen, S.; Yang, J.; Wang, J.; Yao, L.; Yao, X.; Cui, P.; Xu, X. Interface ReEngineering of Li10GeP2S12 Electrolyte and Lithium Anode for All-Solid-State Lithium Batteries with Ultralong Cycle Life. ACS Appl. Mater. Interfaces 2018, 10, 2556–2565. (144) Oh, G.; Hirayama, M.; Kwon, O.; Suzuki, K.; Kanno, R. Bulk-Type All Solid-State Batteries with 5 V Class LiNi0.5Mn1.5O4 Cathode and Li10GeP2S12 Solid Electrolyte. Chem. Mater. 2016, 28, 2634–2640. (145) Sakuda, A.; Hayashi, A.; Tatsumisago, M. Interfacial Observation between LiCoO2 Electrode and Li2S−P2S5 Solid Electrolytes of All-Solid-State Lithium Secondary Batteries Using Transmission Electron Microscopy. Chem. Mater. 2010, 22, 949–956. (146) Auvergniot, J.; Cassel, A.; Ledeuil, J. B.; Viallet, V.; Seznec, V.; Dedryvère, R. Interface Stability of Argyrodite Li6PS5Cl toward LiCoO2, LiNi1/3Co1/3Mn1/3O2, and LiM2O4 in Bulk All-Solid-State Batteries. Chem. Mater. 2017, 29, 3883–3890. (147) Hakari, T.; Deguchi, M.; Mitsuhara, K.; Ohta, T.; Saito, K.; Orikasa, Y.; Uchimoto, Y.; Kowada, Y.; Hayashi, A.; Tatsumisago, M. Structural and Electronic-State Changes of a Sulfide Solid Electrolyte during the Li Deinsertion-Insertion Processes. Chem. Mater. 2017, 29, 4768–4774. (148) Haruyama, J.; Sodeyama, K.; Tateyama, Y. Cation Mixing Properties toward Co Diffusion at the LiCoO2 Cathode/Sulfide Electrolyte Interface in a Solid-State Battery. ACS Appl. Mater. Interfaces 2017, 9, 286–292. (149) Otoyama, M.; Ito, Y.; Hayashi, A.; Tatsumisago, M. Raman Imaging for LiCoO2 Composite Positive Electrodes in All-Solid-State Lithium Batteries Using Li2S-P2S5 Solid Electrolytes. J. Power Sources 2016, 302, 419–425. (150) Zhang, W.; Leichtweiß, T.; Culver, S. P.; Koerver, R.; Das, D.; Weber, D. A.; Zeier, W. G.; Janek, J. The Detrimental Effects of Carbon Additives in Li10GeP2S12-Based Solid-State Batteries. ACS Appl. Mater. Int. 2017, 9, 35888–35896. (151) Auvergniot, J.; Cassel, A.; Foix, D.; Viallet, V.; Seznec, V.; Dedryvère, R. Redox Activity of Argyrodite Li6PS5Cl Electrolyte in All-Solid-State Li-Ion Battery: An XPS Study. Solid State Ionics 2017, 300, 78–85. (152) Mo, Y.; Ong, S. P.; Ceder, G. First Principles Study of the Li10GeP2S12 Lithium Super Ionic Conductor Material. Chem. Mater. 2012, 24, 15–17. (153) Dietrich, C.; Weber, D. A.; Culver, S.; Senyshyn, A.; Sedlmaier, S. J.; Indris, S.; Janek, J.; Zeier, W. G. Synthesis, Structural Characterization, and Lithium Ion Conductivity of the Lithium Thiophosphate Li2P2S6. Inorg. Chem. 2017, 56, 6681–6687. (154) Yamane, H.; Shibata, M.; Shimane, Y.; Junke, T.; Seino, Y.; Adams, S.; Minami, K.; Hayashi, A.; Tatsumisago, M. Crystal Structure of a Superionic Conductor, Li7P3S11. Solid State Ionics 2007, 178, 1163–1167. 33 ACS Paragon Plus Environment


(155) Han, F.; Gao, T.; Zhu, Y.; Gaskell, K. J.; Wang, C. A Battery Made from a Single Material. Adv. Mater. 2015, 27, 3473–3483. (156) Sumita, M.; Tanaka, Y.; Ohno, T. Possible Polymerization of PS4 at a Li3PS4/FePO4 Interface with Reduction of the FePO4 Phase. J. Phys. Chem. C 2017, 121, 9698–9704. (157) Hakari, T.; Deguchi, M.; Mitsuhara, K.; Ohta, T.; Saito, K.; Orikasa, Y.; Uchimoto, Y.; Kowada, Y.; Hayashi, A.; Tatsumisago, M. Structural and Electronic State Changes of a Sulfide Solid Electrolyte during the Li Deinsertion/Insertion Processes. Chem. Mater. 2017, 4768–4774. (158) Yu, C.; Ganapathy, S.; Eck, E. R. H. V.; Wang, H.; Basak, S.; Li, Z.; Wagemaker, M. Accessing the Bottleneck in All-Solid State Batteries, Lithium-Ion Transport over the Solid-Electrolyte-Electrode Interface. Nat. Commun. 2017, 8, 1–9. (159) Sumita, M.; Tanaka, Y.; Ikeda, M.; Ohno, T. Charged and Discharged States of Cathode/Sulfide Electrolyte Interfaces in All-Solid-State Lithium Ion Batteries. J. Phys. Chem. C 2016, 120, 13332–13339. (160) Strauss, F.; Bartsch, T.; de Biasi, L.; Kim, A.-Y.; Janek, J.; Hartmann, P.; Brezesinski, T. Impact of Cathode Material Particle Size on the Capacity of Bulk-Type All-SolidState Batteries. ACS Energy Lett. 2018, 3, 992–996. (161) Nam, Y. J.; Oh, D. Y.; Jung, S. H.; Jung, Y. S. Toward Practical All-Solid-State Lithium-Ion Batteries with High Energy Density and Safety: Comparative Study for Electrodes Fabricated by Dry- and Slurry-Mixing Processes. J. Power Sources 2018, 375, 93–101. (162) Oh, D. Y.; Kim, D. H.; Jung, S. H.; Han, J.-G.; Choi, N.-S.; Jung, Y. S. Single-Step Wet-Chemical Fabrication of Sheet-Type Electrodes from Solid-Electrolyte Precursors for All-Solid-State Lithium-Ion Batteries. J. Mater. Chem. A 2017, 5, 20771–20779. (163) Yao, X.; Liu, D.; Wang, C.; Long, P.; Peng, G.; Hu, Y. S.; Li, H.; Chen, L.; Xu, X. High-Energy All-Solid-State Lithium Batteries with Ultralong Cycle Life. Nano Lett. 2016, 16, 7148–7154. (164) Yang, Y.; Wu, Q.; Cui, Y.; Chen, Y.; Shi, S.; Wang, R. Z.; Yan, H. Elastic Properties, Defect Thermodynamics, Electrochemical Window, Phase Stability, and Li+ Mobility of Li3PS4: Insights from First-Principles Calculations. ACS Appl. Mater. Interfaces 2016, 8, 25229–25242. (165) Pugh, S. F. XCII. Relations between the Elastic Moduli and the Plastic Properties of Polycrystalline Pure Metals. London, Edinburgh, Dublin Philos. Mag. J. Sci. 1954, 45, 823–843. (166) Kondrakov, A. O.; Schmidt, A.; Xu, J.; Geßwein, H.; Mönig, R.; Hartmann, P.; Sommer, H.; Brezesinski, T.; Janek, J. On the Anisotropic Lattice Strain and Mechanical Degradation of High- and Low-Nickel NCM Cathode Materials for Li-Ion Batteries. J. Phys. Chem. C 2017, 121, 3286–3294. (167) Reimers, J. N.; Dahn, J. R. Electrochemical and In Situ X-Ray Diffraction Studies of Lithium Intercalation in LixCoO2. J. Electrochem. Soc. 1992, 139, 2091–2097. (168) Schweidler, S.; de Biasi, L.; Schiele, A.; Hartmann, P.; Brezesinski, T.; Janek, J. Volume Changes of Graphite Anodes Revisited: A Combined Operando X-Ray Diffraction and In Situ Pressure Analysis Study. J. Phys. Chem. C 2018, 122, 8829– 8835. (169) Machida, N.; Kashiwagi, J.; Naito, M.; Shigematsu, T. Electrochemical Properties of All-Solid-State Batteries with ZrO2-Coated LiNi1/3Mn1/3Co1/3O2 as Cathode Materials. Solid State Ionics 2012, 225, 354–358. (170) Sakuda, A.; Takeuchi, T.; Kobayashi, H. Electrode Morphology in All-Solid-State Lithium Secondary Batteries Consisting of LiNi1/3Co1/3Mn1/3O2 and Li2S-P2S5 Solid Electrolytes. Solid State Ionics 2016, 285, 112–117. 34 ACS Paragon Plus Environment

Page 34 of 35

Page 35 of 35 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


(171) Tian, H.-K.; Qi, Y. Simulation of the Effect of Contact Area Loss in All-Solid-State LiIon Batteries. J. Electrochem. Soc. 2017, 164, E3512–E3521.

For table of contents only


Designing ionic conductors: the interplay between structural

Recommend Documents