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Cite This: ACS Appl. Mater. Interfaces XXXX, XXX, XXX−XXX

Atomistic Assessments of Lithium-Ion Conduction Behavior in Glass−Ceramic Lithium Thiophosphates Ji-Su Kim,† Wo Dum Jung,† Ji-Won Son,† Jong-Ho Lee,† Byung-Kook Kim,† Kyung-Yoon Chung,‡ Hun-Gi Jung,‡ and Hyoungchul Kim*,† †

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High-Temperature Energy Materials Research Center, Korea Institute of Science and Technology, 5 Hwarang-ro 14-gil, Seongbuk-gu, Seoul 02792, Republic of Korea ‡ Center for Energy Storage Research, Korea Institute of Science and Technology, 5 Hwarang-ro 14-gil, Seongbuk-gu, Seoul 02792, Republic of Korea S Supporting Information *

ABSTRACT: We determined the interatomic potentials of the Li-[PS43−] building block in (Li2S)0.75(P2S5)0.25 (LPS) and predicted the Li-ion conductivity (σLi) of glass−ceramic LPS from molecular dynamics. The Liion conduction characteristics in the crystalline/interfacial/glassy structure were decomposed by considering the structural ordering differences. The superior σLi of the glassy LPS could be attributed to the fact that ∼40% of its structure consists of the short-ranged cubic S-sublattice instead of the hexagonally close-packed γ-phase. This glassy LPS has a σLi of 4.08 × 10−1 mS cm−1, an improvement of ∼100 times relative to that of the γ-phase, which is in agreement with the experiments.

KEYWORDS: glass−ceramic, molecular dynamics, interatomic potentials, sulfide electrolyte, lithium thiophosphate

U

these materials. Although a variety of Li-ion superconducting crystalline SEs with Li-ion conductivity (σLi) of 10 mS cm−1 or more such as Li10GeP3S1223 and Li9.54Si1.74P1.44S11.7Cl0.324 were recently reported, the many advantages of glassy materials have continued to be the driving force behind further research in glass-based SEs.10−14 In particular, the discovery of glass− ceramic structures reported over a decade ago is an important achievement that doubles the potential of glassy structures.12,13 The σLi value of glass−ceramic SEs obtained by formation of the highly Li-ion conducting Li7P3S11 metastable crystal structure in the glassy matrix was 1 order of magnitude higher than that of existing glassy material.12−14 These results of glass−ceramic SEs are important for achieving the same high σLi as crystalline SEs while maintaining the merits of glassy SEs.10−14,20,21 However, the lack of long-range ordering in the glass structure is a major disadvantage when attempting computational or experimental analyses, and it is difficult to apply a small unit-cell structure with periodic boundary conditions or perform diffraction analysis.25 Furthermore, because a glass−ceramic material is a composite structure of glassy and crystalline phases, Li-ion migration in these materials has not been understood until now, neither has a glass−ceramic structure been realized in practice.

nlike crystalline solids, which have an ordered structure consisting of repeating unit cells, amorphous solids are composed of randomly arranged atoms or molecules and have crystallographic characteristics that lack long-range ordering. Thus, amorphous materials have short-range order and varying degrees at a few atomic or molecular dimensions.1−5 The crystallographic characteristics of these amorphous solids convey unique properties (e.g., isotropic,6 gradually softening over a temperature range,7 and an irregular shape and surface8,9) and these materials are widely used in various applications. In recent years, glassy materials and their derivatives have been studied as Li-ion conductive solid electrolytes (SEs) for all-solid-state battery (ASSB) applications.10−15 Glass-based materials used as SEs for ASSBs vary widely from oxides to sulfides.10−19 In the early days of research, the development of materials focused on those based on conventional oxides such as silicate16,17 and phosphate.18,19 However, subsequently, research and development progressed to the sulfide-based SEs analogs of the oxide-based glasses, which were known to have superior material properties (e.g., high ionic conductivity,10−14 low processing cost,20,21 excellent formability22) two decades ago. A representative form of a sulfide-based SE is a combination of Li2S and P2S5 as glass modifier and network former, respectively.1,10,14 Most sulfidebased SEs reported thus far, regardless as to whether they are crystalline or amorphous, are derived from a combination of © XXXX American Chemical Society

Received: October 8, 2018 Accepted: December 21, 2018

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DOI: 10.1021/acsami.8b17524 ACS Appl. Mater. Interfaces XXXX, XXX, XXX−XXX

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ACS Applied Materials & Interfaces

Table 1. Interatomic Potential Parameters for Crystalline γLi3PS4a

This Letter presents an analysis of the Li-ion migration behavior of glass−ceramic (Li2S)0.75(P2S5)0.25 (LPS) structures, a topic that has not yet been investigated, by employing molecular dynamics (MD) simulations. We investigate and predict the differences in the S-sublattice and σLi values determined by different crystallographic structures (in this work, glass, the crystalline γ-Li3PS4, and glass−ceramic in fixed LPS composition), and we compare them to the available experimental results. The effect of the contributions of each phase in the glass−ceramic structure is also addressed by using local region analysis. As mentioned above, an amorphous solid is a material with crystallographic characteristics that lack long-range ordering. However, an amorphous material also has a short-range internal structure consisting of interconnected structural blocks. Sulfide-based Li-ion conductive SEs have the thiophosphate PS43− anion as a basic building block, as is the case for commonly seen oxide-based glasses, silicate and phosphate materials consisting of SiO43− and PO43− clusters, respectively. Generally, for the LPS structure, we confirmed that the PS43− tetrahedron, the basic building block, is repeated, although it lacks long-range ordering. Obviously, there are various anion clusters such as P2S62− and P2S74−, depending on the oxidation states and coordination numbers of P, but the PS43− cluster is dominant in an LPS composition consisting of Li2S and P2S5 in a molar ratio of 0.75:0.25, irrespective of whether the structure is amorphous or crystalline. We focused on this feature of amorphous materials for our comprehensive computational studies of LPS materials in glass or glass−ceramic structures, which is the objective of this research. Specifically, we implemented empirical potential fields to simulate the interactions of PS43− tetrahedrons with Li-ions using the thermomechanical properties (e.g., the lattice parameters and elastic moduli) of the γ-Li3PS4 crystal structure (see Figure S1) obtained from the aforementioned LPS ratio composition. To reasonably represent the ionic and covalent bonds that constitute Li-[PS43−] in γ-Li3PS4, we combine the well-known long-range Coulombic, short-range Morse, and three-body harmonic potentials. The two-body interaction EI of the Coulombic and Morse potentials was expressed as EI =

qiqj εr

e 2 + D[e−2α(r − r°) − 2e−2α(r − r°)]

potential two-body, EI

potential three-body, EII

atom i

atom j

Li Li P S atom i S

Li S S S

D

0.058000000 0.040755682 0.410420000 0.024096411 atom j atom k P

S

α

r0◦ (Å)

3.9870 1.3988 2.3287 1.3585 K (eV)

3.40378 3.20378 2.20028 4.28352 θ◦ (deg)

0.4

108.50

The effective atomic charges of Li, P, and S are +0.88,+ 1.28, −0.98, respectively. a

temperature−pressure (NPT) ensemble as the following process. (i) Because the LPS glass structure has the same PS43− tetrahedral anion clusters as mentioned above, the ionic and Coulombic interactions of the glassy LPS structure can be determined with the Li-[PS43−] potential of the γ-structure. (ii) We constructed a glassy LPS with a randomly arranged atomic structure without long-range ordering by relaxing a low-density cell (with sufficiently large lattice parameters; see Figure 1a) with the NPT ensemble to obtain a stable structure (Figure 1b). We observed no notable structural change with the variation of the initial annealing temperature (see Figure S2). Figure 1 shows the process of obtaining this glassy structure and shows that the glassy LPS structure obtained after equilibrium has very reasonable Li and PS43− distributions that lack long-range ordering (Figure 1c). The radialdistribution function (RDF) associated with the glassy LPS structure thus obtained is shown in Figure S3 of the Supporting Information, respectively. As shown in Figure 2, the glass−ceramic LPS structure was constructed by combining the optimized γ-Li3PS4 unit cell and glassy LPS structure. The lattice mismatch between the γLi3PS4 unit cell and glass LPS structure was reduced by constructing a 10.78 Å × 10.78 Å × 10.78 Å cell-sized glassy LPS structure. Then, the equilibrated glass structure was used to construct the glass (210) surface with a surface size of 43.12 Å × 24.16 Å. The γ-Li3PS4 (100) surface was constructed with a surface size of 39.27 Å × 24.18 Å. By combination of the glassy LPS (210) and γ-Li3PS4 (100) surfaces, the glass− ceramic structure was constructed with a cell size of 39.25 Å × 24.18 Å × 114 Å. After the equilibration process of the glass− ceramic structure was performed, the final cell size and density of the glass−ceramic LPS structure were 40.58 Å × 22.68 Å × 107.46 Å and 1.95 g cm−3, respectively. The final glass− ceramic structure thus obtained had both glassy and crystalline phases, and an interfacial region with an approximate width of 20 Å was observed in the middle region (see the inset of Figure 2b). In the interfacial region, the evolution of the atomic arrangement from crystalline to glass or vice versa was observed, and detailed atomic distributions and RDF changes are shown in Figures S4 and S5, respectively. The density (ρ) values obtained from each part of the glass−ceramic LPS structure are shown in Figure 2c as a function of temperature. These results agree well with the reported computational and experimental results, and we conclude that our glass−ceramic structure constructed here is reasonable. To analyze the Li-ion transport characteristics of the glass− ceramic LPS structure, we collected the Li-ion trajectories for each region for a period of 500 ps with an MD simulation. Details of the calculation are listed in the Supporting Information. Figure 3a shows an atomistic illustration of the

(1)

where qi and qj are the ionic charges of the ions i and j, e is the elementary charge, ε is the dielectric constant, r is the distance between the ions i and j, r◦ is the equilibrium distance, D is the well-depth, and α is the stiffness parameter, respectively. The three-body harmonic potential EII was expressed as E II = K(θ − θ )2 (2) ° where K is the prefactor and θ◦ is the equilibrium angle, respectively. Each potential parameter is matched to the reported thermomechanical properties of γ-Li3PS4 by using density-functional theory and potential optimization techniques (see the Supporting Information, further computational details). The final converged fitting values are summarized in Table 1. For your reference, the thermomechanical properties of γ-Li3PS4 obtained from the final potential fields are also listed in Table S1 and are in good agreement with the available experimental results. The LPS glass structure was implemented with the interatomic potential fields of the γ-Li3PS4 and the constantB

DOI: 10.1021/acsami.8b17524 ACS Appl. Mater. Interfaces XXXX, XXX, XXX−XXX

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ACS Applied Materials & Interfaces

Figure 1. Generation of glass LPS structure with Li-ions and PS43− anion clusters. Images of the atomic structure of the glass LPS supercell structure (a) before and (b) after equilibration. (c) Distribution of Li-ions (top) and PS43− anion clusters (bottom) in the glass LPS structure. The white, blue, and yellow spheres indicate the Li-, P-, and S-ions, respectively.

Figure 2. Generation of glass−ceramic LPS structure with glass LPS and γ-Li3PS4 structures. (a) Two initial structures were combined to construct (b) the final glass−ceramic structure. The inset is a magnified image to show the atomic structure in detail in the vicinity of the interface (i.e., the area indicated by the red cuboid). (c) Variations of the predicted LPS density in the glass−ceramic structure as a function of temperature. The available computational26 and experimental26,27 results are also listed.

Figure 3. (a) Trajectories of Li-ions in glass−ceramic LPS structure at T = 650 K for 500 ps. The red line indicates the Li-ion trajectory. (b) Averaged MSD for crystal, interface, and glass area with calculation time. (c) Arrhenius plot of predicted σLi of glass−ceramic LPS structure using MD simulations. Experimental results28,29 of typical LPS sulfides including crystalline γ-Li3PS4 and glass phases are also presented. For your reference, the related Li-ion diffusivity results are shown in Figure S6.

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DOI: 10.1021/acsami.8b17524 ACS Appl. Mater. Interfaces XXXX, XXX, XXX−XXX

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ACS Applied Materials & Interfaces

Table 2. Reported and Calculated Ea and σLi of LPS Glass and Related Crystalline Phases at Room Temperature (25 °C)a LPS phase

compound

method

Ea (eV)

glass interface crystal crystal crystal glass

(Li2S)0.75(P2S5)0.25 (Li2S)0.75(P2S5)0.25 γ-Li3PS4 γ-Li3PS4 β-Li3PS4 (Li2S)0.75(P2S5)0.25

MD MD MD Exp. AIMD AIMD

0.28 0.29 0.36 0.49 0.35 0.17

σLi (mS cm−1) 4.08 7.07 4.35 3.00 8.79 9.88

× × × × × ×

10−1 10−2 10−3 10−4 10−1 10−2

note this work this work this work ref 29 ref 30 ref 26

a

AIMD stands for ab-initio molecular dynamics.

Li-ion trajectories obtained for each region at 650 K and clearly shows the difference in the Li-ion trajectories induced by the structural difference between regions. In the crystalline γ-phase, Li-ions are constrained near the apex of the PS43− tetrahedrons, whereas Li-ions are actively moving across the tetrahedrons in the glassy and interfacial structures. On the basis of the results obtained for the mean-square displacement (MSD), which is a quantitative value of the Li-ion movement, the MSD value of the interfacial region was approximately 2 times as large as that of the crystalline phase, whereas the MSD of the glassy matter was approximately 6 times as large as that of the crystalline structure. On the basis of the MSD values, we predicted the value of σLi using the following relationship. σLi =

c Li(qLie)2 kBT

DLi

(3)

where kB is the Boltzmann constant, cLi is the Li-ion concentration, qLi is the Li-ion charge state, and DLi is the Li-ion diffusion coefficient over time t [=MSD(t)/6t]. Figure 3c is the Arrhenius plot of the change in σLi with temperature calculated by MD simulations. To verify the reliability of the predicted σLi of the glass−ceramic structures, the σLi of various LPS materials is also shown (see Figure 3c and Table 2).26,28−30 As the MSD results show, the highest σLi of 4.08 × 10−1 mS cm−1 was observed for the glassy phase, whereas the lowest σLi was 4.35 × 10−3 mS cm−1 in the crystalline phase. In the interfacial region, the value of σLi, which is less than that of the glassy structure because of the suppression of Li-ion migration through the crystalline region, is approximately 7.07 × 10−2 mS cm−1 at room temperature. As reported recently, the migration of Li-ions is known to be sensitive to the S-sublattice structure.31,32 According to the literature results, the BCC S-sublattice act as a fast Li-ion pathway composed of the S4 tetrahedron, resulting in a lower Li-ion migration barrier energy. The fact that the value of σLi differs in each region was confirmed for the glass−ceramic structure; however, an S-sublattice analysis of the glassy and interfacial structures has not yet been reported because of the absence of long-range ordering. In this work, we developed a method to analyze the local regions. This method divides the MD structure into subsections and examines the S-sublattice for each local region. Figure 4a shows the local S-sublattice distribution obtained from the glass−ceramic structure. As is well-known, the crystalline γ-phase has a hexagonally closepacked (HCP) S-sublattice, whereas the glassy and interfacial structure is a mixture of HCP and cubic S-sublattice.31,32 As shown in Figure S7 for RDF analysis, the longer and broader features of the inter S−S bonds in glassy S-sublattice indicate the formation of the local BCC and disorder structures, respectively. Figure 4b, which shows the variations of the Ssublattice ratio in the z-axis direction, clearly shows this change

Figure 4. S-sublattice distribution in the glass−ceramic LPS structure. (a) Visual representation of S-sublattice distribution assigned by local area analysis. The pink, green, and gray indicate the HCP, cubic, and other S-sublattices, respectively. (b) Variation of each S-sublattice ratio along the z-axis.

in the S-sublattice with different structures. The cubic Ssublattice grows rapidly in the interfacial region where the long-range ordering of the crystalline γ-phase collapses and increases to constitute approximately 40% of the structure. Locally, the cubic S-sublattice is not continuously distributed, but this ratio remains constant in the glassy region, ensuring that the σLi of the glass remains high. Thus, the cubic Ssublattice ratio of approximately 40% in the interfacial and glassy regions results in an activation energy (Ea) of approximately 0.28 eV (see Table 2), which is very low compared with that of crystalline γ-phase. From the viewpoint of structural dynamics, as clarified, the glassy LPS structure (i.e., Li and PS43− distributions) lacking long-range ordering led to significant atomic disorder and large displacement D

DOI: 10.1021/acsami.8b17524 ACS Appl. Mater. Interfaces XXXX, XXX, XXX−XXX

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ACS Applied Materials & Interfaces (Figures 1c and S8). Considering that the Li-ion conduction is dynamically coupled to the Li-ion site disorder and PS43− anion reorientation,30,33 we reasonably expect a high σLi in the glassy structure. Our results confirmed that the arrangement and reorientation of anion clusters greatly affects the values of σLi even for the same composition ratio. We also believe that the cubic S-sublattice found in the local domain can provide new insights that would be useful to improve the σLi of the PS43− cluster-based SEs. Although we identified the features of the fast Li-ion conducting structure of sulfide-glass in this study, it is difficult to experimentally control the local structure of the glass structure. However, considering that the glass structure consists of building blocks composed of a glass modifier and a network former under high energy conditions, we expect the tailoring of the local glass structure to enable the tuning of its Li-ion conduction by employing various control techniques (e.g., heat treatment conditions, element coordination number). In addition, a well-known glass−ceramic structure with high σLi is a (Li2S)0.70(P2S5)0.30 structure, Li7P3S11. Unfortunately, this study focuses only on structures based on the interatomic potentials of the Li-[PS43−] building block, so we cannot proceed with the (Li2S)0.70(P2S5)0.30 glass−ceramic analysis. In the near future, we expect to analyze the local structure and ion conduction phenomena of (Li2S)0.70(P2S5)0.30 glass−ceramic structures through the development of Li-[P2S74−] potential and the construction of Li7P3S11 structure. In summary, we conducted a comprehensive analysis of the Li-ion transport characteristics of the glass−ceramic LPS structure using MD simulations and examined the evolution of the S-sublattice in the interfacial and glassy domains for improved Li-ion conduction. First, we identified unique PS43− anion clusters in the short-range region of the glassy structure and newly developed the interactive force fields of the Li[PS43−] bonds. Employing these empirical potentials, we succeeded in constructing the glass−ceramic structure (including both the crystalline γ-Li3PS4 and interfacial and glassy LPS structures) and used these results to conduct MD simulations. Using the Li-ion trajectories obtained from the MD simulations, we predicted the σLi of each region, and these values agree with the experimental results. By introducing the local S-sublattice analysis, we also found that 40% of the glassy and interfacial region comprised cubic S-sublattice, unlike the crystalline γ-phase, which consisted of only HCP. These Ssublattice evolutions in the interfacial and glassy domains led to a significant reduction of Ea and notable improvement of σLi. These results provide new directions for understanding glass− ceramic sulfides with the view of achieving superior Li-ion conducting materials as an SE for ASSBs.

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Crystal data, LPS glass-ceramic structure (CIF)

AUTHOR INFORMATION

Corresponding Author

*Hyoungchul Kim. E-mail: [email protected]. ORCID

Ji-Won Son: 0000-0002-5310-0633 Jong-Ho Lee: 0000-0003-4481-6258 Kyung-Yoon Chung: 0000-0002-1273-746X Hun-Gi Jung: 0000-0002-2162-2680 Hyoungchul Kim: 0000-0003-3109-660X Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS This work was supported by the Energy Efficiency & Resources Core Technology Program of the Korea Institute of Energy Technology Evaluation and Planning (KETEP) via financial resources from the Ministry of Trade, Industry & Energy, Republic of Korea (No. 20152020106100). This work was partly supported by the Dual Use Technology Program of the Institute of Civil Military Technology Cooperation via financial resources from the Ministry of Trade, Industry & Energy and Defense Acquisition Program Administration (17CM-EN-11). This research was partly supported by the Technology Development Program to Solve Climate Changes of the National Research Foundation (NRF) funded by the Ministry of Science and ICT (grant number: 2017M1A2A2044482).

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ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acsami.8b17524. Further computational details; structural information and thermomechanical properties of γ-Li3PS4; equilibrium process of glass LPS structure; the RDF and MSD analyses of glass and glass−ceramic LPS structures; Arrhenius plot of Li-ion diffusivity in the glass−ceramic LPS structure (PDF) Crystal data, LPS glass structure (CIF) E

DOI: 10.1021/acsami.8b17524 ACS Appl. Mater. Interfaces XXXX, XXX, XXX−XXX

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DOI: 10.1021/acsami.8b17524 ACS Appl. Mater. Interfaces XXXX, XXX, XXX−XXX