Fast Lithium Ion Migration in Room Temperature LiBH4 - The Journal

Article pubs.acs.org/JPCC

Fast Lithium Ion Migration in Room Temperature LiBH4 Young-Su Lee* and Young Whan Cho High Temperature Energy Materials Research Center, Korea Institute of Science and Technology, Seoul 02792, Republic of Korea S Supporting Information *

ABSTRACT: The defect structure and the Li ion diffusion mechanism of orthorhombic LiBH4 (o-LiBH4) are studied by first-principles calculations to elucidate the Li ion transport in o-LiBH4. Two metastable Li interstitial sites are identified, and the formation energies of the Schottky and Frenkel defect pair are calculated to be 1.2−1.4 eV, the former being slightly easier to form. The energy required to form intrinsic defects is higher than that of hexagonal LiBH4 (h-LiBH4). On the other hand, the migration energy barrier of the Li vacancy or interstitial ranges from 0.1 to 0.3 eV, which is comparable to that of hLiBH4. Therefore, the higher defect formation energy mainly accounts for the much lower Li ion conductivity of o-LiBH4. The calculated overall activation barrier for the Li ion conduction is in fair agreement with the experimental activation energy. Molecular dynamics simulation demonstrates that both the interstitial and the interstitialcy mechanisms are operative for the Li interstitial diffusion and that the interconnected interstitial sites compose a fast diffusion path. The simulation results point out that the enhancement of the carrier density via defect or interface engineering may significantly raise the ionic conductivity of o-LiBH4.

1. INTRODUCTION As battery technology becomes a key player in the zero emission vehicles, interest in viable solid electrolytes has increased to replace current liquid electrolytes in the lithium ion batteries in pursuit of improved safety and performance.1 Among several types of candidate solid electrolytes, metal borohydrides and their derivatives stand as a relatively new kind, but recent years have seen a rapid development in this field,2−4 continuously adding new materials, such as LiBH4,5 Li2B12H12,6,7 LiCe(BH4)3Cl,8 Na3BH4B12H12,9 NaCB9H10,10 and Na2B10H10,11 to the list of promising candidates for a lithium or sodium ion conductor. This type of materials is well-known for structural diversity, and sudden increase in ionic conductivity is often associated with a polymorphic transformation as are the cases for LiBH4,5 NaCB9H10,10 LiNaB12H12,7 etc. The most thoroughly studied case is LiBH4 as it initiated the exploration of metal borohydrides as an ionic conductor. The ionic conductivity of the high temperature (> 110 °C) polymorph in the hexagonal structure (h-LiBH4) is over 10−3 S cm−1,12 and the feasibility of making an all-solid-state Li ion battery with h-LiBH4 as an electrolyte was tested.13,14 The ionic conductivity of room temperature orthorhombic phase (o-LiBH4), however, is about 10−8 S cm−1, far below being practical. Interestingly, the room temperature conductivity of LiBH4 can be dramatically increased to 10−4 S cm−1 by infiltrating it into mesoporous scaffolds or mixing with nanoparticles.15,16 Such infiltrated LiBH4 was employed as an electrolyte in the all-solid-state electrochemical cell.17 The enhanced Li ion diffusion at the interface with C, SiO2, or Al2O3 has been demonstrated, mostly by NMR spectroscopy,15,18−22 but its atomistic origin is still © 2017 American Chemical Society

illusive. Unless the crystal structure of o-LiBH4 is completely destroyed at the interface, we may be able to interpret the increased interface conductivity in terms of the defect structure of o-LiBH4. Conductivity rise at the interface can be originated either from increased carrier density or from enhanced mobility since the conductivity is the product of the two terms. Carrier density increase by a defect accumulation in the space charge regions at the interface has long been discussed,23−25 and in this sense, understanding the defect structure is indispensable. There have been both experimental and theoretical studies to elucidate the defect structure of highly conducting h-LiBH4. The activation energy for the Li ion conduction is 0.53−0.56 eV,5,21 and the migration energy barrier has been estimated to be 0.31 eV by first-principles calculations.26 The discrepancy between the experimental activation energy and the calculated migration barrier can be attributed to the formation energy of the intrinsic defects. For example, in LiBH4:LiI solid solution, the activation energy and the calculated migration energy are 0.63−0.68 eV27 and 0.2−0.3 eV,28 respectively, and the calculated Frenkel pair formation energy of 0.44 eV28 roughly fills the gap. Relatively little attention has been given to the intrinsic defects and the Li ion diffusion in o-LiBH4. Earlier investigation on the native defects in o-LiBH429 was focused on the defects related to the hydrogen storage properties not to the Li ion transport. To provide a deeper understanding on the Li ion transport in LiBH4, we herein report the structure and the formation energy of the intrinsic defects related to the Li Received: June 28, 2017 Revised: August 3, 2017 Published: August 4, 2017 17773

DOI: 10.1021/acs.jpcc.7b06328 J. Phys. Chem. C 2017, 121, 17773−17779

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The Journal of Physical Chemistry C ion diffusion and propose diffusion mechanisms from firstprinciples calculations.

2. CALCULATION METHODS The defect formation energies and the migration barriers for Li ions were investigated by first-principles calculations. The calculations were performed within the framework of density functionaly theory using the generalized-gradient approximation by Perdew, Burke, and Ernzerhof30 for exchangecorrelation functional, implemented in the Vienna Ab-initio Simulation Package.31,32 Projector augmented wave potential33 with a plane-wave cutoff energy of 600 eV was used. Atomic coordinates were optimized until the force on each atom became smaller than 0.005 eV/Å. For the simulation of a Frenkel or a Schottky pair, a 2 × 3 × 2 supercell containing 48 formula units (288 atoms) of o-LiBH4 was used and even a larger 3 × 5 × 3 supercell (1080 atoms) was employed in order to test the convergence. Only the single Γ-point was sampled for efficiency. The minimum energy path of Li ion migration was obtained by the climbing-image nudged elastic band method (CI-NEB).34 In addition to the zero temperature estimation of the migration energy barrier, finite temperature dynamics of Li ions was simulated. Born−Oppenheimer molecular dynamics simulation was carried out in a canonical ensemble (NVT) with the Nosé−Hoover thermostat fixed at 450 K. The simulation temperature was higher than room temperature to speed up the Li ion diffusion. The system was equilibrated for 10 ps, and the simulation ran for 70 ps with a time step of 1 fs.

Figure 1. Formation energies (Ef) of the Schottky and Frenkel defect pairs as a function of the distance (r) between the pair. The inset is the same plot as a function of 1/r. The dashed lines in the inset are a guide to the eye.

more favorable, but they are likely to coexist since the difference is minor. Notably, the Frenkel pair shows stronger dependence on 1/r because the interstitial charge is more localized and less effectively shielded by the nearby ions. For comparison, a linear line with the slope of −e2/4πε is drawn in the inset, where e is unit charge and ε is permittivity of LiBH4.35 This slope represents the Coulomb energy variation between two unit charges, +e and −e, embedded in the dielectric medium of LiBH4 and is similar to the variation of the Frenkel pair formation energy. The thermodynamic equilibrium in eqs 1a and 1b gives the intrinsic defect concentration proportional to exp(−Ef/2kT), where k and T denote the Boltzmann constant and the absolute temperature, respectively. The defect fraction at room temperature is about 10−11 for Ef = 1.3 eV, which is too small to deliver a high conductivity unless the mobility is extremely high. On the other hand, the calculated Ef of a Frenkel pair in h-LiBH4 is as low as 0.44 eV,28 giving an equilibrium defect fraction of 2 × 10−4. The reported formation energy might have been underestimated since the interstitial and the vacancy are closely located in their study,28 but even the closest Frenkel pair in our study requires 0.88 eV to form, confirming the higher defect formation energy in o-LiBH4. Therefore, it is clear that a lesser number of mobile Li ions are available in o-LiBH4 compared to h-LiBH4, which explains the much lower conductivity of oLiBH4. Now we elaborate on the structural change upon introducing a Frenkel defect. We identified two metastable interstitial sites: one is 4c Wyckoff position and the other is 4b position in the Pnma space group36 as listed in Table 1. The 4c interstitial site is more stable by ∼0.1 eV (see the formation energy in Figure 1), and its geometry is illustrated in Figure 2. The 4c site features a coordination by four [BH4]− groups: one tridentate, two bidentate, and one monodentate groups. It is definitely less tightly bound compared to the lattice Li site coordinated by 9 H atoms, but still has 8 coordinating H atoms. The coordination environment is summarized in Table S1, and the 4c interstitial sites viewed in different directions are presented in Figure S1. Intriguingly, we find a structural analogy with h-LiBH4. The metastable interstitial site in h-LiBH4 in the P63mc space

3. RESULTS AND DISCUSSION 3.1. Defect Formation. The formation of the two most common types of defect pairs in ionic crystals, i.e., Schottky and Frenkel pairs, is considered. The defect formation reactions in the Kröger−Vink notation are as follows: • ′ + V BH null ↔ V Li 4

′ + Lii• LiLi× ↔ V Li

Ef = ΔE

Ef = ΔE

(1a) (1b)

One pair of defects is embedded in the simulation cell containing 288 atoms (48 f.u. of LiBH4), and the distance (r) between the two defect sites, V′Li−V•BH4 for a Schottky pair and V′Li−Li•i for a Frenkel pair, is varied within the simulation cell. The defect formation energies (Ef) as a function of the aforementioned distance are presented in Figure 1. The formation energy was calculated by subtracting the energy of the equiatomic bulk LiBH4 from the energy of the simulation cell having one pair of defects. No further correction was applied. The error coming from the periodic boundary condition would be small since we treat neutral paired defects which preserve the original stoichiometry. In the case of a Frenkel defect, ca. 4 Å is the shortest V′Li−Li•i distance at which a defect pair can remain stable: the vacancy and the interstitial spontaneously recombine and the structure then becomes defect-free when the distance is shorter. Due to the Coulomb interaction, the formation energy is linearly dependent on 1/r, which is shown in the inset of Figure 1. To check the convergence, a single calculation was performed using an even larger supercell containing 1080 atoms (180 f.u. of LiBH4) and the limit of r→∞ (y intercept in the inset) is estimated by an extrapolation. The formation energies thus obtained are approximately 1.2−1.4 eV, with a Schottky pair being slightly 17774


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The Journal of Physical Chemistry C

interconnected Li, B, and the interstitial sites in o-LiBH4 form a corrugated layer, whereas those in h-LiBH4 make up an almost flat layer as illustrated in Figure 3. The structural similarity hints that the polymorphic transformation would proceed by reorientation and/or displacement of the constituting ions without long-range diffusion.37 More importantly, such minor rearrangement results in the significant change in the defect formation energy, i.e., strongly affects the stability of the interstitial sites. In this respect, the reorganization of the ions at the interface38 may account for the enhanced interface conductivity observed in some LiBH4 composites as it can lower the defect formation energy, thereby drastically increasing the carrier density. 3.2. Defect Migration. We proceed further to the discussion on the migration of a vacancy and an interstitial. First, Li ion migration energy barriers to the nearest vacant positions were assessed. The four nearest sites considered, V1, V2, V3, and V4, are drawn as polyhedrons with the coordinating [BH4]− groups as vertices in Figure 4. The Li

Table 1. Calculated Li, B, H, and Two Metastable Li Interstitial Sites in o-LiBH4 (Pnma Space Group)a Li B H1 H2 H3 Li•i Li•i a

Wyckoff symbol

x

y

z

4c 4c 4c 4c 8d 4c 4b

0.1546 0.3167 0.9128 0.4090 0.2181 0.122 1/2

1/4 1/4 1/4 1/4 0.0251 1/4 1/2

0.1123 0.4192 0.9282 0.2627 0.4175 0.713 0

The cell parameters are a = 7.3746 Å, b = 4.3819 Å, and c = 6.5409 Å.

Figure 2. Coordination environment of the lattice Li sites (green) and the 4c interstitial sites (red) is drawn in a polyhedron model, with coordinating H atoms at the vertices.

group36 is surrounded by three Li ions and three B atoms as shown in Figure 3b.26 The crystal structure of o-LiBH4 viewed along the a direction (see Figure 3a) has a similarly looking hexagonal arrangement,26,36,37 and the 4c interstitial site also has three Li ions and B atoms nearby. The difference is that the

Figure 4. Li ion migration to the four nearest vacant sites listed in Table 2. The Li sites and the 4c interstitial sites are displayed as green and red polyhedrons, respectively, with B atoms at the vertices.

site shares an edge with the nearest V1 site and a corner with the other three sites. The minimum energy paths from the CINEB calculation are presented in Figure 5a. The calculation predicts that the most favorable migration path is a jump to the third nearest V3 site, not to the closest V1. The low migration energy barrier to V3 can be attributed to the locally stable intermediate state around ξ ∼ 0.5 in the energy profile in Figure 5a. We found that the moving Li ion is located close to the 4c interstitial site at this local minimum point, and so is the case with V4. Therefore, the interstitial site plays a vital role in suppressing the energy barrier as a stable intermediate position for a vacancy migration. The atomic configurations are presented in Figure S2. Second, the migration of an interstitial Li ion from a Frenkel pair was analyzed. Different from a vacancy migration through the lattice sites, both interstitial and interstitialcy mechanisms39 can be operative in the Li interstitial migration. The energy profiles along the two most favorable migration paths via the interstitial (I) and the interstitialcy (IC) mechanisms are plotted in Figure 5b. The energy barriers are in general even lower than those for a vacancy migration in Figure 5a, suggesting that the Li interstitials would diffuse faster than the Li vacancies. The lowest energy barrier for each migration mechanism is I2 (0.107 eV) < IC2 (0.245 eV) < V3 (0.297 eV).

Figure 3. Bonding environment around the Li interstitial site (red) for (a) o-LiBH4 (from our calculation) and (b) h-LiBH4 (from ref 35). The distances (in Å) between the Li interstitial to the Li site (green) and to the B site (gray) are annotated in the upper panel. The lower panel illustrates the atomic arrangement viewed along a different direction, with the intralayer thickness and the interlayer distance annotated. 17775


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Figure 5. Minimum energy path for (a) a vacancy migration and (b) an interstitial migration.

network of the 4c sites bridged by the 4b sites extended along the b direction and is visualized in Figure 6b. In the interstitialcy mechanism, an interstitial Li ion (Li1) knocks off a nearby Li ion at the lattice site (Li2) and that Li ion (Li2) then becomes an interstitial. To gain more insight on the interstitialcy mechanism, we plot in Figure 7 the normalized

Table 2. Hopping Distances and Associated Energy Barriers for a Defect Migration via the Vacancy (V), Interstitial (I), and Interstitialcy (IC) Mechanismsa V1 V2 V3 V4 I1 I2 IC1 IC2

distance (Å)

energy (eV)

3.487 4.104 4.180 4.382 3.720 3.968 4.368 (2.626, 2.820) 4.708 (2.626, 3.204)

0.310 0.383 0.297 0.360 0.320 0.107 0.262 0.245

a

In the case of the interstitialcy mechanism, the hopping distances of the two Li ions involved are written inside the parentheses.

The energy barrier of I2 is remarkably low, and such low activation barrier again relies on the stable intermediate state at ξ ∼ 0.5 in Figure 5b. The atomic configuration of the stable intermediate state (see Figure 6a) displays the Li interstitial coordinated by six H atoms; the Li interstitial is located at the 4b interstitial site in Table 1. The I2 migration path builds a

Figure 7. Normalized distance to the final position for the initial (Li1) and the final (Li2) interstitial ion along the minimum energy path of IC1 and IC2 in Figure 5b.

distance to the final position for the two Li ions involved (Li1 and Li2). The variations in atomic positions demonstrate distinctly different behavior in IC1 and IC2. In the case of IC1, Li2 escapes first from the original site, and then there are two Li interstitials sitting nearby at the energy maximum point (ξ ∼ 0.5). Finally, Li1 takes the position of Li2 releasing the energy. On the other hand, almost simultaneous displacement of Li1 and Li2 toward the final position takes place in the case of IC2. The atomic configurations are presented in Figure S3. Overall, the defect migration energy barriers from the CINEB calculations are 0.1−0.3 eV, and these values are similar to 0.2−0.3 eV of h-LiBH4.26,28 Therefore, the dramatic increase in conductivity upon phase transformation from o-LiBH4 to hLiBH4 is likely due to an increase in defect concentration rather than due to an increase in mobility. Notably, the predicted migration barriers of o-LiBH4 are very low and even comparable to the activation energies of some other borohydrides such as

Figure 6. (a) The atomic configuration at the stable intermediate state of I2 in Figure 5b. The yellow atom is the moving interstitial ion at the 4b site, and the red ones are the initial and the final 4c interstitial positions. (b) The predicted lowest energy barrier path for the migration of Li interstitials highlighted in red ellipsoids. The 4c-4b-4c interstitial sites are connected as drawn in red lines, and the network is extended along the b direction. 17776


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The Journal of Physical Chemistry C Li2B12H12 (0.34 eV)40 and LiM(BH4)3Cl (0.3−0.4 eV)41,42 that exhibit very high room temperature ionic conductivity. Combining the result of the defect formation energy (Ef) and the migration energy (Em), we estimate the overall activation energy (Ea) of the Li ion conduction in o-LiBH4. Since the conductivity is the product of the carrier density and the mobility, the activation energy is the sum of the two terms, Ea = Ef/2 + Em. The activation energy thus derived is 0.9−1.0 eV for the vacancy and 0.75−0.9 eV for the interstitial. This value is in fair agreement with the experimental activation energy of 0.7− 0.8 eV5,14 obtained from the Arrhenius plot of the ionic conductivity of o-LiBH4, which validates our interpretation on the defect formation and migration. We would like to emphasize that the Li ion conductivity of o-LiBH4 can become significantly high if enough numbers of carriers are created via defect engineering43 or interface engineering.25 3.3. Molecular Dynamics Simulation. The motion of Li ions at finite temperature was simulated by molecular dynamics (MD). A Frenkel or Schottky pair, or an interstitial Li ion, was introduced in the 288 atom simulation cell. In the case of the paired defects, the two defect sites were placed at the farthest distance allowed in the given simulation cell at the beginning. However, as the simulation time progressed, the Li vacancy either recombined with the Li interstitial or was bound to the [BH4]− vacancy, and Li migration did not proceed afterward. We therefore focus on the simulation cell containing an extra Li ion which is bound to remain as an interstitial. We note in passing that o-LiBH4 does not display any significant disorder, different from the case of h-LiBH4.44−46 Nonetheless, the thermal motion of Li and B is anisotropic and pronounced along the a direction (see Figure S4), in agreement with the experimental atomic displacement parameters at 360 K.45 In addition, the enhanced thermal motion slightly displaces Li and B so that the corrugated layer composed of Li and B in Figure 3a becomes thinner, reflecting the gradual transition toward hLiBH4 in Figure 3b. Another interesting observation is that the isosurfaces of the H1 population are separated from those of H2 and H3, and at the same time more localized, while the isosurfaces of H2 and H3 are more closely connected and spread out (see Figure S4). This difference in thermal behavior is again consistent with the experimental atomic displacement parameters of H atoms45 and indicates that the libration or rotation of [BH4]− around the B-H1 axis would be dominant. Overall, the important features found through the experimental structure analysis are reproduced in our MD simulation. By analyzing the trajectory of the interstitial, we found that both interstitialcy and interstitial mechanisms are operative for the interstitial diffusion. The variation in the x, y, and z coordinates (fractional coordinates in the a, b, and c crystal axis, respectively) of the Li ions in Figure 8a captures the exchange between the interstitial and an adjacent lattice Li ion. The background color is changed whenever the exchange takes place. As visualized in Figure 8a, the Li interstitial sometimes travels quite a long distance (interstitial mechanism) before finally settling down at a lattice site (interstitialcy mechanism), but there are occasions where several exchange events occur almost simultaneously. To find out any directionality in diffusion, the mean square displacement in the a, b, and c directions is calculated and plotted in Figure 8b. The result indicates relatively facile diffusion along the b direction, although the diffusion is not strictly contained in one dimension. The faster diffusion along the b direction corroborates the migration through the 4c-4b-4c

Figure 8. (a) Trajectory of the moving Li ions. Background color change indicates an exchange occasion between the interstitial and a lattice Li ion. (b) Mean square displacement and its decomposition into the a, b, and c directions.

interstitial network (I2 mechanism). The Li ion population density provides a more direct picture on the preferred Li diffusion path. The isosurface in Figure 9a reveals the increased Li ion density in the interstitial region (marked as a yellow circle), which again coincides with the highlighted region in Figure 6b. We note that the simulation time is not long enough to get an even distribution over symmetrically equivalent sites. Figure 9b is a cross section of the plane of the lattice Li ions as marked in Figure 9d. The increased Li population density appears as a trail from the Li site to the 4c interstitial site; the Li interstitial seems to linger for a while before or after replacing the lattice Li ion. The cross sections in Figure 9c,d show increased population around the 4b interstitial site (Figure 9c) and along the b direction (Figure 9d), highlighting the role of the 4b interstitial site and the most energetically favorable diffusion path. Therefore, both the migration barrier at 0 K and the ion dynamics at 450 K consistently predict the fast diffusion channel in o-LiBH4 composed by the metastable interstitial sites.

4. CONCLUSIONS Our first-principles study on the defect structure and the Li ion migration in o-LiBH4 shows that the much lower Li ion conductivity of o-LiBH4 compared to h-LiBH4 is due to higher defect formation energy rather than due to lower mobility. In 17777


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Figure 9. (a) Isosurface of the Li ion population density. Cross sections (b) through the lattice Li ions and (c, d) through the 4b interstitial sites in different view directions. The positions of the cross sections are displayed as white dotted lines in (a) and (d). Li ion density contour lines are drawn on each cross-sectional image, with increasing density from blue to red. Rechargeable Batteries. Appl. Phys. A: Mater. Sci. Process. 2016, 122, 251. (4) Paskevicius, M.; Jepsen, L. H.; Schouwink, P.; Cerny, R.; Ravnsbaek, D. B.; Filinchuk, Y.; Dornheim, M.; Besenbacher, F.; Jensen, T. R. Metal Borohydrides and Derivatives - Synthesis, Structure and Properties. Chem. Soc. Rev. 2017, 46, 1565−1634. (5) Matsuo, M.; Nakamori, Y.; Orimo, S.; Maekawa, H.; Takamura, H. Lithium Superionic Conduction in Lithium Borohydride Accompanied by Structural Transition. Appl. Phys. Lett. 2007, 91, 224103. (6) Her, J.-H.; Yousufuddin, M.; Zhou, W.; Jalisatgi, S. S.; Kulleck, J. G.; Zan, J. A.; Hwang, S.-J.; Bowman, R. C.; Udovic, T. J. Crystal Structure of Li2B12H12: A Possible Intermediate Species in the Decomposition of LiBH4. Inorg. Chem. 2008, 47, 9757−9759. (7) He, L.; Li, H.-W.; Nakajima, H.; Tumanov, N.; Filinchuk, Y.; Hwang, S.-J.; Sharma, M.; Hagemann, H.; Akiba, E. Synthesis of a Bimetallic Dodecaborate LiNaB12H12 with Outstanding Superionic Conductivity. Chem. Mater. 2015, 27, 5483−5486. (8) Ley, M. B.; Ravnsbæk, D. B.; Filinchuk, Y.; Lee, Y.-S.; Janot, R.; Cho, Y. W.; Skibsted, J.; Jensen, T. R. LiCe(BH4)3Cl, a New LithiumIon Conductor and Hydrogen Storage Material with Isolated Tetranuclear Anionic Clusters. Chem. Mater. 2012, 24, 1654−1663. (9) Sadikin, Y.; Brighi, M.; Schouwink, P.; Č erný, R. Superionic Conduction of Sodium and Lithium in Anion-Mixed Hydroborates Na3BH4B12H12 and (Li0.7Na0.3)3BH4B12H12. Adv. Energy Mater. 2015, 5, 1501016. (10) Tang, W. S.; Matsuo, M.; Wu, H.; Stavila, V.; Zhou, W.; Talin, A. A.; Soloninin, A. V.; Skoryunov, R. V.; Babanova, O. A.; Skripov, A. V.; et al. Liquid-Like Ionic Conduction in Solid Lithium and Sodium Monocarba-Closo-Decaborates near or at Room Temperature. Adv. Energy Mater. 2016, 6, 1502237. (11) Udovic, T. J.; Matsuo, M.; Tang, W. S.; Wu, H.; Stavila, V.; Soloninin, A. V.; Skoryunov, R. V.; Babanova, O. A.; Skripov, A. V.; Rush, J. J.; et al. Exceptional Superionic Conductivity in Disordered Sodium Decahydro-Closo-Decaborate. Adv. Mater. 2014, 26, 7622− 7626. (12) Maekawa, H.; Matsuo, M.; Takamura, H.; Ando, M.; Noda, Y.; Karahashi, T.; Orimo, S. I. Halide-Stabilized LiBH4, a RoomTemperature Lithium Fast-Ion Conductor. J. Am. Chem. Soc. 2009, 131, 894−895. (13) Unemoto, A.; Yasaku, S.; Nogami, G.; Tazawa, M.; Taniguchi, M.; Matsuo, M.; Ikeshoji, T.; Orimo, S.-i. Development of Bulk-Type All-Solid-State Lithium-Sulfur Battery Using LiBH4 Electrolyte. Appl. Phys. Lett. 2014, 105, 083901. (14) Sveinbjörnsson, D.; Christiansen, A. S.; Viskinde, R.; Norby, P.; Vegge, T. The LiBH4-LiI Solid Solution as an Electrolyte in an AllSolid-State Battery. J. Electrochem. Soc. 2014, 161, A1432−A1439. (15) Blanchard, D.; Nale, A.; Sveinbjörnsson, D.; Eggenhuisen, T. M.; Verkuijlen, M. H. W.; Suwarno; Vegge, T.; Kentgens, A. P. M.; de Jongh, P. E. Nanoconfined LiBH4 as a Fast Lithium Ion Conductor. Adv. Funct. Mater. 2015, 25, 184−192. (16) Choi, Y. S.; Lee, Y.-S.; Oh, K. H.; Cho, Y. W. InterfaceEnhanced Li Ion Conduction in a LiBH4-SiO2 Solid Electrolyte. Phys. Chem. Chem. Phys. 2016, 18, 22540−22547.

addition, the structural analogy between o-LiBH4 and h-LiBH4 suggests that even a minor reorganization of the Li and [BH4]− ions can significantly lower the defect formation energy, which is likely to be present at the interface. Therefore, the enhanced interface conductivity found in the LiBH4-SiO2 composites can be attributed to the increased defect concentration either by lowered formation energy or by defect accumulation at the space charge region since the mobility is predicted to be sufficiently high. The molecular dynamics simulation demonstrates that the interconnected metastable interstitial sites serve as the lowest energy path for the interstitial diffusion. The facile Li ion migration in o-LiBH4 revealed in our study shows a promise for a significant conductivity enhancement in o-LiBH4 by increasing the carrier density via defect or interface engineering.

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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/acs.jpcc.7b06328. Table S1 and Figures S1−S4 (PDF)

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AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]. Tel: +82-2-958-5412. ORCID

Young-Su Lee: 0000-0002-3160-6633 Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS This work was supported by the Convergence Agenda Program (CAP) of the Korea Research Council of Fundamental Science and Technology (KRCF) (Grant number CAP-11-05-KRISS) and by the Innovation Fund Denmark via the research project HyFill-Fast.

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REFERENCES

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