Modified Anion Packing of Na2B12H12 in Close to Room

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Modified Anion Packing of Na2B12H12 in Close to Room Temperature Superionic Conductors Yolanda Sadikin,† Pascal Schouwink,†,# Matteo Brighi,† Zbigniew Łodziana,‡ and Radovan Č erný*,† †

Department of Quantum Matter Physics, Laboratory of Crystallography, University of Geneva, Quai Ernest-Ansermet 24, CH-1211 Geneva, Switzerland ‡ Polish Academy of Sciences, Institute of Nuclear Physics, ul. Radzikowskiego 152, 31-342 Kraków, Poland S Supporting Information *

ABSTRACT: Three different types of anion packing, i.e., hexagonal close packed (hcp), cubic close packed (ccp), and body centered cubic (bcc), are investigated experimentally and with ab initio calculations in the model system Na2B12H12. Solvent free and water assisted mechanical grinding provide polycrystalline samples for temperaturedependent synchrotron radiation X-ray powder diffraction and electrochemical impedance spectroscopy. It is shown that among the common close packed lattices, the hcp anionic backbone creates very favorable conditions for three-dimensional ionic conduction pathways, comprised of O−O, T−T, and T−O−T (O for octahedral, T for tetrahedral) cation hops. The hcp lattice is stable with respect to ccp and bcc lattices only at higher volumes per formula unit, which is achieved either by cationic substitution with larger cations or partial substitution of hydrogen by iodine on the closoanion. It is found that the partial cationic substitution of sodium with lithium, potassium, or cesium does not lead to enhanced conductivity due to the obstruction of the conduction pathway by the larger cation located on the octahedral site. Substitution on the closo-anion itself shows remarkable positive effects, the ionic conductivity of Na2B12H12‑xIx reaching values of close to 10−1 S cm−1 at a rather low temperature of 360 K. While the absolute value of σ is comparable to that of NaCB11H12, the temperature at which it is attained is approximately 20 K lower. The activation energy of 140 meV is determined from the Arrhenius relation and among the lowest ever reported for a Na-conducting solid.

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INTRODUCTION Large-scale battery grids for electrochemical energy storage are becoming one of the principal prospective solutions in the development of modulated electricity supply systems aimed at storing energy harvested from solar wind or hydro-installations in the form of chemical energy and delivering on demand. One of the ways to eliminate safety issues in secondary batteries is going “all solid” in the future. This transition has begun for Liion batteries, and daily progress is improving the manufacturability and stability of solid-state based devices for state of the art mobile technologies. Na-Based Battery Systems. All solid state sodium batteries are an enticing prospect for large-scale grid storage, due to both the geographical distribution and abundance of Na as compared to Li metal, and the softer restrictions imposed on the energy density (i.e., energy per unit volume or weight) in stationary applications. While the electrochemical potential of Na/Na+ is just slightly lower than that of Li/Li+ (2.71 vs 3.04 V vs SHE), the real strength of Na-assemblies needs to lie in cheap, robust, and long-lived large-scale installations, where energy density and output voltage are not ultimately crucial. Such solutions have been realized in the form of Na−S and ZEBRA cells, which use β″-alumina as a solid electrolyte, and operate at high temperatures (ht) near 573 K (300 °C), where both electrodes are molten. Heating to these temperatures is © 2017 American Chemical Society

costly and entails additional safety issues in the case of Na−S technologies. By replacing β″-alumina with NASICON it has been possible to lower operating temperatures down to 400 K in some cases.1 Currently, extensive efforts are concentrated on designing solid electrolytes capable of conducting ions at roomtemperature (rt), which exhibit a sufficiently high Naconductivity above 10−3 S/cm and a low activation energy.2−4 Next to eliminating heating costs, rt operation implies that the capacity may be tremendously increased with respect to graphitic and other systems by the usage of Na metal in the solid state (melting point 371 K). A solid material minimizes flammability issues and battery failure due to solvent reactions, but entails its own challenges arising due to the solid nature of the material. Some of these, the mechanical properties with respect to thermal expansion, the operating voltage window, and the chemical reactivity at the electrolyte-electrode interface, have been found to be well addressed by the consideration of complex hydride electrolytes.5−8 The complex hydride LiBH4 initially triggered materials research in this field as a solid hydrogen store, the focus, however, currently shifting to the higher metal boranes, capable of conducting both Li, Na, and even Mg (the latter in solution). Received: January 9, 2017 Published: April 11, 2017 5006

DOI: 10.1021/acs.inorgchem.7b00013 Inorg. Chem. 2017, 56, 5006−5016

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Inorganic Chemistry

Structure−Property Design of Metal-Hydroborates. Recent reports15−17 point toward the chemical modification of the [B12H12]2− anion as means of structural engineering. The resulting anions [CB11H12]− and [CB9H10]− are monovalent and stabilize two types of materials, LiCB11H12 and NaCB11H12, and LiCB9H10 and NaCB9H10, respectively. While in the conducting dodeca-crystal structures the anion sublattice corresponds to a ccp configuration, which stands in contrast to the bcc concept, the anion packing in the conducting decacrystal structures is of hcp type. The hcp anion sublattice is conserved also in the very recently reported conducting anionmixed phases Li2(CB9H10)(CB11H12) and Na2(CB9H10)(CB11H12).18 As is the case with closo-boranes, the high charge carrier mobility requires activation through polymorphic order−disorder transformations, though the energy barriers are modest due to the single negative charge and hence a lower “sticking” factor of the mobile cation to the anion lattice. Yet, it is remarkable that transition temperatures are shifted to below 400 K in homoleptic carboranes and even below 300 K in Na2(CB9H10)(CB11H12), with respect to critical temperatures of 529 and 615 K for the closo-borane materials of Na and Li, respectively.19 Very recently, the above-mentioned closocarborane anion [CB11H12]− was successfully employed in the development of a halogen-free liquid electrolyte for Mgbatteries,20 a milestone in its own. Modification of the borane anion is possible also by substituting on the hydrogen site. Some effort has been invested by the hydrogen storage community to replace hydrogen by fluorine in borohydrides, with the aim of optimizing hydrogen release reactions.21 Regarding ionic conductivity, the perchlorated closo-borane [B12Cl12]2− was used as molecular species in nonaqueous electrolytes (US patent US4849310A) as early as 1980.22 Halogenation of the borane anion therefore appears as a promising path to set out on a search for further solid borane electrolytes. In our group we tend to focus on crystal structure modification by replacement schemes outside the individual molecule, substituting entire anions or cations. A cationsubstituted closo-borane LiNaB12H12 with high ionic conductivity of close to 0.8 S cm−1 at 550 K was recently reported as a Na-substituted Li2B12H12 crystallizing in the anti-CaF2 structure type of ccp anion topology, which becomes superionic at ht.23 Again, the disorder-driven conduction mechanism is comparable to that of Na2B12H12, while the conducting phase ht-LiNaB12H12 is still close packed and differs from the bcc lattice of the Na-endmember. The premise we set out with in this study was to explore the surely existing, but hitherto not discovered, hcp borane materials containing Li+ or Na+ cations. We were motivated by the report on hcp boranes for the divalent alkaline earths Ca2+, Sr2+, and Ba2+.24,25 The hcp occurs very frequently among inorganic compounds and could possibly, in the absence of high temperature-induced disorder, show improved mobility and potentially lower activation energies with respect to ccp materials, due to the differences in connectivity between tetrahedral and octahedral interstitials available for cation occupation. In the first part of this article, we show how to create hcpcloso-borane lattices by cation-substitution in mixed-cation compounds. We synthetically investigate a considerable compositional space in order to identify the structural parameters leading to polymorphism and discuss the impact of lattice distortions on the available migration pathways.

Hydroborate-Based Na-Electrolytes. Few very simple pseudobinary metal boranes have proven quite favorable properties, in particular, the dodeca- and decahydroborates (and their derivatives) built from the polyanions [B12H12]2− and [B10H10]2−. Higher metal boranes are ionic materials where the borane-anion is charge-balanced by a metal cation.9,10 The resulting crystal structures are packed, and all common packing variants can be found, the body centered cubic (bcc), cubic close packed (ccp), as well as the hexagonal close packed (hcp), which this article primarily deals with. The tumbling motions of the polyanion commonly lead to different kinds of rotational disorder. The influence of the motions on the anion packing in the crystal can result in enhancing the ionic mobility as well as lowering the activation energy for conduction pathways in the structure. For instance, the high temperature polymorphs of Na2B12H12 (bcc) and Na2B10H10 (ccp) exhibit ionic conductivities of 0.1 and 0.01 S cm−1 at 540 and 383 K, respectively, with activation energies of 210 and 470 meV.11,12 In both cases, however, the superionic ht phase is limited to temperatures above the melting point of Na metal and is not stable at rt. The conduction in these phases is activated by entropy, implying significant structural dynamics of the borane anion and resulting in slight, but significant, modifications of the anion packing. Bearing in mind that technological considerations call for high ionic mobilities close to rt, we recently managed to stabilize a high conductivity of close to 10−3 S cm−1 at 298 K, by a targeted manipulation of the anion lattice of a borane material, consisting of anion-mixing of the large closo-borane [B 12 H 1 2 ] 2 − and the smaller borohydride anion in Na3BH4B12H12.13 This value is comparable to that of NASICON-type materials or β-alumina. We found that substituting on the anion and cation sites simultaneously may lead to a rich phase diagram of double-cation solid electrolytes, which do not require rotational disorder in order to conduct, and which we described the first member of, (Li0.7Na0.3)3BH4B12H12.13 In contrast to the more simple metal boranes, the conduction mechanism in these mixedanion boranes is driven by vacancies rather than molecular reorientations. (Li0.7Na0.3)3BH4B12H12, however, is not stable below 473 K, since the chemical reaction forming it at 498 K is reversible upon cooling. Nevertheless, this suggests an approach to a purposeful and structured search for novel borane conductors. A further general design concept to obtain high ionic conductivities is believed to reside in the necessity of building a bcc-packed anion lattice serving as a backbone to the ionic conductor. Such a packing contains tetrahedral voids which percolate by sharing faces and in which the mobile species are loosely bounded. This idea dates back to the first investigations on the superionic bcc phases of AgI or Ag2S. The concept was recently revisited very instructively by means of ab initio calculations on sulfide materials, currently the best crystalline solid state Li-conductors, Li 10 GeP 2 S 12 and Li7P3S11.14 In line with this assumption, the first solid state borane electrolyte reported to exhibit superionic mobility, the ht-phase of Na2B12H12 has an anionic bcc backbone. However, the bcc arrangement of anions does not occur very frequently among inorganic materials as the tetrahedral and octahedral voids do not have regular shape and further design principles need to be identified for these large anion hydroborate salts. This calls for an investigation of the topological effects in the much more frequently occurring close-packed topologies and the analysis of accessible interstitials, in the context of higher boranes. 5007

DOI: 10.1021/acs.inorgchem.7b00013 Inorg. Chem. 2017, 56, 5006−5016

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both O- and T-sites present low energy environments for a mobile species, the hcp may therefore be more favorable to cationic disorder, given that three different types of hops; i.e., O−O, T−T, and T−O are possible. Polymorphism of Na- and LiAB12H12 (A = Li, Na, K, Cs). The previously reported polymorphs of Li2B12H12, Na2B12H12, and the solid solution LiNaB12H12 crystallize in densely packed ccp anion lattices.11,19,23 By changing the cation size-ratio in mixed-metal boranes, we modify the anion backbone and obtain highly stable hcp lattices. Hereby, the n O-interstitials are fully occupied by the larger cation, while the smaller one fills n T-interstitials, leading to an occupancy of the tetrahedral site of 50%. These results are summarized in Figure 1. The lattice

Additionally, the phase diagram of Na2B12H12 is rigorously reinvestigated and extended, resolving mysteries recently observed in the polymorphism of this material. In the second part of this article, we report on a novel hcppacked Na-closo-borane material Na2B12H12‑xIx, where the anion substitution H− ↔ I− stabilizes the hcp lattice, which is host to the remarkable Na-mobility of 4 × 10−2 S cm−1 at T > 362 K (below the melting point of Na metal) at an ultralow, liquid-like activation energy of Ea = 140 meV, even below that of the recently reported Na2Fe2(SO4)3.26 These values set a new benchmark in the approach to rt superionic behavior, with respect to metal boranes as well as the most competitive NASICON-type materials at the respective temperature, the latter requiring special postsynthetic treatment to attain comparable values. We place this finding into context with the topology of distorted hcp borane lattices and their impact on structural bottlenecks in the migration pathways of mobile species. This packing variant has previously not been reported, and we describe how the hcp-lattice is capable of opening bottlenecks for ionic conduction. In the third part, the structural parameters stabilizing hcp configurations are corroborated with extensive density functional theory (DFT) calculations, and ionic diffusion is investigated theoretically in the bcc, ccp, and hcp lattices with molecular dynamics (MD). A conduction mechanism is suggested that does not necessarily require disorder in ht-phases, but much rather resembles that of an interstitially driven electrolyte.

Figure 1. Anion packing (in color code) and polymorphism of metal closo-boranes.

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RESULTS AND DISCUSSION On the basis of the large number of ionic compounds displaying superionic behavior in packed crystal structures, it seems conceptually sensible to search for close-packed ionic conductors built from bulky quasi-spherical borane anions. It can be expected that the resulting materials will be simple in structure, allow for easy chemical modification, and are related to their precursors due to the dominating influence of the anion-size in the design of the crystal packing. We systematically screened this binary compositional materials space by performing mechanochemistry on equimolar mixtures of alkalimetal closo-boranes. In all cases simple addition reactions are observed yielding phases of composition NaAB12H12 in cationordered compounds, and three solid solutions, (Li,K)2B12H12, (Li,Cs)2B12H12 and (Na,K)2B12H12. With the exception of the two cubic polymorphs of (Na,K)2B12H12, all the discovered compounds and polymorphs reported on herein belong to the structure aristotype Ni2In. There exist three different packing variants relevant to the following discussion. The packing density of a bcc-packing of hard spheres amounts to 68% of space, while the total packing fraction of the densely packed ccp and hcp is identical (∼74%). For a net of n atoms the bcc lattice has 3n octahedral and 6n tetrahedral interstitials (Figure S1). Again, the ccp and hcp lattices have an identical number of interstitials, n octahedral (O) and 2n tetrahedral (T), the principal difference between them being the distribution and the connectivity between the interstitials, in particular, regarding out-of-plane-connecting channels. In ccp, the Osites share faces with eight adjacent T-sites, eliminating potential O−O hops in ccp solid state ionic conductors. The hcp configuration does allow such moves, given the face-sharing O-channels parallel to the stacking direction (crystallographic caxis). The T-voids alternatively share faces and edges along the same direction and form trigonal bipyramids. The in-plane connectivity is identical in both stacking variants. Provided that

parameters are listed exhaustively in Table 1. Difference plots from Rietveld refinement and fits from pattern modeling as well as temperature-dependent synchrotron radiation X-ray powder diffraction (SR-XPD) data are given in Supporting Information as Figures S2−S16. We identify a critical size ratio at the Li−K and Na−K boundary (Figure 1), suggesting that it is the ionic size of the octahedrally coordinated K+ ion that induces the hcp. The phase diagram of (Na,K)2B12H12 contains all three packing variants (Figure 1). This points toward rather small energetic differences between packing variants, which appears counterintuitive given the differences in density. We will come back to this when discussing the phase diagram of Na2B12H12 below. Surprisingly, it was possible to obtain satisfying Rietveld fits during refinement for all novel compounds without significant orientational disorder of the anion. This is a relevant observation, as it points out conduction mechanisms that do not need to be “switched on” by order−disorder transformations; rt conductivity may therefore become feasible. NaCsB12H12. The temperature-dependent structural evolution of hcp packed metal closo-boranes (green in Figure 1) is best exemplified on the mixed-cation compound NaCsB12H12. The observations and conclusions obtained on this compound can likewise be transferred to the related materials. This material is triclinic at room temperature and increases symmetry stepwise to monoclinic, orthorhombic, and finally the hexagonal parent phase. It is both intriguing and suspicious that a symmetry as low as P1̅ should be stable in such simple packed ionic materials, and we emphasize that the diffraction pattern cannot be indexed in the monoclinic crystal system. The presence of four different polymorphs of NaCsB12H12 allows us to follow the structural evolution of the underlying hcp anion lattice as a function of temperature as it approaches the ideal geometry with a c/a axis ratio of 1.633. In order to 5008

DOI: 10.1021/acs.inorgchem.7b00013 Inorg. Chem. 2017, 56, 5006−5016

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Table 1. Space Group (SG) Symmetry and Lattice Parameters Determined at Temperature T of Mixed Metal closo-Boranes Reported in This Work (t.w.) compound

SG

V [Å3]

a [Å]

b [Å]

c [Å]

o-LiCsB12H12 h-LiCsB12H12 o-(Li,Cs)2B12H12 c-(Li,Na)2B12H12

Fm3̅ Pm3̅n

12.5 7.25 9.42 9.83 10.13 7.19 7.58 7.62 12.455 7.53 7.20 7.49 7.31 10.65 8.30 8.42

12.05 12.13 10.43

c1-(Na,K)2B12H12 c2-(Na,K)2B12H12

1079.98 552.97 1340.14 949.01 1040.20 536.79 1119.50 1138.66 1188.66 602.46 527.30 581.55 538.66 1207.95 572.82 598.01

7.17

h-(Li,K)2B12H12 tri-NaCsB12H12 m-NaCsB12H12 o-NaCsB12H12 h-NaCsB12H12 tri-(Na,K)2B12H12 h-(Na,K)2B12H12

Pn21a P63mc Cmc21 Pa3̅ Fm3̅m P63mc P1̅ P21/c Pn21a P63mc P1 P63mc

13.64

11.99 12.27 12.29 12.44 12.26 11.68 11.97 11.64

12.06 12.17 7.66 7.21

α [°]

β [°]

γ [°]

91.54

91.77 92.35

91.29

92.69

87.11

119.40

T [K]

ref

300 308 297 303 513 657 325 434 675 747 300 773 303 308 547 755

t.w. t.w. t.w. 23, 27 23 t.w. t.w. t.w. t.w. t.w. t.w. t.w. t.w. t.w.

temperature increase, these displacements are successively lost until only one mode is active in orthorhombic NaCsB12H12, which corresponds solely to shifts of A and B planes against each other. This mode vanishes in the hexagonal phase. While the ideal hcp is approached with temperature, the c/a ratio changes from 1.62−1.61−1.63−1.63, no longer changing in the ortho-hex due to one sole in-plane distortion being active in the orthorhombic phase (Figure 2). It is important to note that all Rietveld refinements were tested for mixed Na/Cs occupancies, allowing us to exclude such disorder with the precision of 0.01 for the occupation factor. In the structural description of hcp, NaCsB12H12 n Osites (100%) are hence occupied by Cs while 50% of the T-sites are occupied in an off-center position by Na cations. This cation distribution will be identical for all hcp-materials presented herein, with the exception of the solid solutions (Li,K)2B12H12, (Li,Cs)2B12H12, and (Na,K)2B12H12. The cation ordering in these lattices will be relevant to the ionic conductivities of these samples. (Na,K)2B12H12. The c1 and c2 phases of (Na,K)2B12H12 (Table 1) have polymorphic ccp-bcc relationships similar to the two nonconducting rt and ht polymorphs of Na2B12H12. The triclinic and hexagonal polymorphs (as well as LiKB12H12) are all derived from Ni2In, and hence comparable to the discussed NaCsB12H12. As all packing variants occur in the (Na,K)2B12H12 phase diagram (no clear explanation has been found concerning the mutual stabilities), this cation combination defines the border between the ccp-packing for smaller cations and the hcp-packing for larger ones on the octahedral site. It is assumed that, in particular in this phase, the differences in the enthalpies are marginal for different anion packing. Phase-Diagram of Na2B12H12. The temperature-induced polymorphic transformations occurring in Na2B12H12 have been studied previously.11,19 The existence of a fourth phase in the diagram was noted, but has evaded further characterization.19 We have revisited this phase diagram and describe this polymorph as a ccp material of space group symmetry Fm3m ̅ . The in situ diffraction data (Figure 3) show the polymorphism as a function of temperature. It can be seen that the monoclinic, both Na- and anion-ordered rt-phase (deformed ccp) splits into a bcc and close-packed branch in a very narrow temperature

visualize this polymorphism in an illustrative way, we have created a centrosymmetric hexagonal parent anion lattice which is a common supergroup (P63/mcm) to all non-hexagonal polymorphs, while conserving the anion positions of hexagonal NaCsB12H12 (P63mc) and have then decomposed the different distortions defining the low-symmetry polymorphs with respect to their symmetry using the program Amplimodes28 available on the Bilbao Crystallographic Server. The resulting distortion modes are shown for each non-hexagonal polymorph of NaCsB12H12 in Figure 2, where the arrows represent the amplitude and direction of the respective distortion of the anion lattice with respect to an ideal hcp of [B12H12]2− anions. In the lowest symmetry triclinic polymorph, four active modes are observed with respect to the parent (Figure 2), two of which result in atomic shifts within the hexagonal plane (P21/m, Pnma) and two perpendicular to it (both P21/c). Upon

Figure 2. Mode decomposition for the lattice distortion in NaCsB12H12 polymorphs. The crystal system of the compound is given at the top, the symmetries of distortion modes on the right. The red spheres represent positions of anions in the ideal hcp; arrows represent the direction and magnitude of atomic shifts due to the respective distortion. 5009

DOI: 10.1021/acs.inorgchem.7b00013 Inorg. Chem. 2017, 56, 5006−5016

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Inorganic Chemistry

succeeded in producing hydrated single crystals of Na2B12H12‑xIx × 4H2O, which very likely contain an identical molecule since the dry closo-variant is produced by dehydration at the boiling point of water under a dynamic vacuum from the hydrated precursor. Crystal Structure of Na2B12H12‑xIx. The crystal structure of the partially halogenated hcp Na-closo-borane is shown in Figure 4, along with the Rietveld refinement plot. As mentioned above, we explain the stabilization of the hcp in this material by iodine substitution on the hydrogen site I− ↔ H−. The fact that minor modification of the anion lattice results in a new compound Na[6]Na[4]B12H12‑xIx corroborates the stability of the Ni2In structure type in these closo-borane salts. The material is synthesized by drying in vacuo from an aqueous methanol solution which results in a fine-grained powder. Unfortunately, neither the iodine content nor atomic position could be determined from powder diffraction data. Nevertheless, single crystals of the precursor (Figure S17), grown by slow evaporation at rt, allowed locating iodine on the specified positions by means of Fourier difference maps. Further evidence of halogenation is provided by energy dispersive X-ray (EDX) maps (Figure S18) on freshly cleaved surfaces of Na2B12H12‑xIx × 4H2O crystals. The hydrated closoborane Na2B12H12 × 4H2O has previously been reported in a tetragonal crystal structure of space group P42/ncm.30 Its symmetry and structural features are conserved upon iodine doping. The electron density map obtained from single crystal diffraction data allowed us to locate four hydrogen positions partially substituted by iodine amounting to a total iodine content of x = 0.1 per closo cluster. While this presents a rather small iodine content, the unit cell expanding merely from 1375.3(2) Å3 to 1376.8(1) Å3 (295 K) upon anion exchange, it is undoubtedly enough to modify the anion-packing from ccp to hcp in the Na-closo compound. The single crystal refinement converges to partial occupancy on those hydride sites which are facing disordered Na-atoms, possibly suggesting some extent of halogen bonding, though this may present only an approximate picture of an average halogenation where many different clososubstitutes are in equilibrium. The crystal structure of Na[6]Na[4]B12H12‑xIx itself is built from an hcp anion-lattice that, upon drying, crystallizes in a monoclinic cell (Pc, a = 7.22, b = 12.54, c = 12.12 Å, β = 90.03°,

Figure 3. Experimentally observed phase evolution of Na2B12H12. The two branches corresponding to different anion lattices are shown in the inset.

range between 513 and 519 K, which then coexist to higher temperatures. Such behavior in an equilibrium state, as we investigate here, is representative of small energy barriers separating different polymorphs. Interestingly, it should be pointed out that the observed densities of ccp and bcc type Na2B12H12 are inversed at 280 and 262 Å3/f.u. (535 K, ΔV/V ≈ 8%), respectively, the bcc density being, unexpectedly, higher than that of the close-packed lattice. The exact same is observed for NaKB12H12 (535 K, ΔV/V ≈ 8%). We will elaborate on these volume-observations in the computational chapter below. During our synthetic efforts to produce closo-boranes by solution chemistry, we discovered a hexagonal phase which was first attributed to a fifth hitherto not described Na2B12H12 polymorph, crystallizing in the same space group symmetry P63mc as the mixed-metal materials discussed above. We then realized that this hexagonal phase is actually stabilized by the anionic substitution I− ↔ H− of terminal hydride moieties of the closo-anion. This substitution readily occurs and proceeds rapidly in aqueous solution.29 It was not possible to locate the iodide within the molecule from powder diffraction since the substitution x in Na2B12H12‑xIx is minor. However, we

Figure 4. (Left) Rietveld plot for data of Na2−I measured at 443 K (λ = 0.81984). Bragg peaks positions from top to the bottom: h-Na2B12H12‑xIx, Na2B12H12−xIx. The refined phase composition is given in wt %. χ2 = 962, Rwp(bgr. corr.) = 0.09. Inset: disordered structure of h-Na2B12H12‑xIx (Na in blue, B in red). (Right) Distortion modes in the Pc rt-phase with respect to the hexagonal ht-polymorph. 5010

DOI: 10.1021/acs.inorgchem.7b00013 Inorg. Chem. 2017, 56, 5006−5016

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Inorganic Chemistry

Figure 5. (Left) Ionic conductivity of a sample containing 52 wt % Na2B12H12‑xIx and 48 wt % NaI. The conductivity of the precursor Na2B12H12 is shown for comparison. (Right) Na-ion conduction channels in hcp sublattice with face-sharing T and O interstices sites in h-Na2B12H12‑xIx. In blue the static Na positions as optimized by DFT.

the new Na closo-borane. We note that the lack of conductivity in hexagonal NaCsB12H12 is owed to the same structural shortcoming, the size of the Cs-cation anchoring it to the octahedral site and preventing potential mobility of a Na-cation loosely located in the T-site. The situation observed in Na2B12H12‑xIx is illustrated in Figure 5 (right). Each octahedron (red) shares six faces with the neighboring T-sites (green) and two faces with octahedra forming channels in the c-direction. The three possible motions between sites are T−T, T−O, and O−O. Conduction pathways were determined using VoronoiDirichlet polyhedra (TOPOS software31) and allowed us to identify the structural bottlenecks for ionic mobility within the hcp lattice (Table S1). The bottleneck size of 1.37 Å encountered in the conducting bcc-phase of Na2B12H1211 was taken as reference value for ionic mobility, this window being large enough to accommodate a Na cation (radius rNa = 1.16 Å). The monoclinic Pc polymorph of Na2B12H12‑xIx has multiple bottlenecks (Table S1) arising due to symmetry lowering and the accompanying lattice distortions (Figure 4 (right)). Several of these have a radius < rNa and reduce the dimensionality of conduction to two-dimensional (2D), defined by the accessible ones with sizes of 1.23−1.77 Å. The theoretically accessible space in this 2D net should provide low energy barriers, which is reflected in rather high and continuously increasing experimental values of the conductivity, already at temperatures below the phase transition (Figure 5, left). In the superionic hexagonal phase the lattice distortions P21/c and Pnma (Figure 4, right) have relaxed and opened up isotropic three-dimensional (3D) conduction with a lower bottleneck size of 1.34 Å. All T−T and O−O hops participate in ionic conduction. It is noteworthy that, despite all hcp- metal closo-boranes sharing the same anion lattice (and seemingly cation positions) only hexagonal Na2B12H12‑xIx hosts an extreme Na-conductivity. The single arc observed in the impedance spectra (Figure S19) is indicative of one single relaxation process. As the measured ionic conductivity of mixed metal hcp-compounds is 2 orders of magnitude lower than that of Na2B12H12‑xIx, it follows that the larger Cs or K cations present obstacles in the conduction pathway, emphasizing the necessity of both T- and O-interstitials being available for the conduction mechanism. The ionic conductivity of Na2B12H12‑xIx reaches values of close to 10−1 S cm−1, comparable to those of NaCB12H12, however, already at a temperature of 360 K, 20 K lower than the carborane. In fact,

T = 300 K, Figure S16), which transforms into the hexagonal phase (P63mc, a = 7.0, c = 12.15, T = 443 K, Figure 4) at approximately 360 K, accompanied by a volume increase of ΔV/V ≈ 3%. The structural data of both phases are given as CIF files in the Supporting Information. During the phase transition the anion positions (Figure 4) approaches ideal values of the hcp, in analogy to the transition sequence discussed for NaCsB12H12. The c/a ratio is very close to that of the ideal hexagonal packing with 1.64, contrary to 1.68 in the monoclinic phase. It is easily seen from Figure 4 that the main distortion amplitude, which is relaxed in the ht-phase, concerns in-plane displacements. In the context of migration pathways this leads to changes in particular of the structural bottlenecks in c-direction, i.e., the O−O hops, which will be discussed below. In Na[6]Na[4]B12H12‑xIx both the octahedral and tetrahedral sites are occupied by the same atom type, the generation of the hcp thus necessarily has its origins in minor anionic modifications. We will discuss the factors changing the anion packing in more detail below. At this stage we point out that experimentally determined unit cell volumes of partially halogenated and hydrogen-pure Na[6]Na[4]B12H12‑xIx differ as much as 1% (569.0 vs 564.1 Å3 at 535 K), the iodination thus resulting in expansion of the unit cell. Rietveld refinement (Figure 4, left) shows that Na[6] occupies a position in the center of the octahedron on average, while the tetrahedral Na[4] is placed off-center, approaching the position of the triangular face shared with the adjacent tetrahedron, leading to the quasitrigonal bipyramidal Na[5] coordination mentioned earlier on. These positions are an indicator of Na-disorder on the T-site. Furthermore, we find no strong experimental evidence for rotational disorder of [B12H12]2−, which has been stated repeatedly to activate superionic behavior in closo-boranes. The “on” switch for ionic conduction may therefore reside in the shifts of hcp layers. The position of Na resembles that of the mobile species in the ccp underlying the conducting phases of LiCB11H12 and NaCB11H12. According to Tang et al.15 the n cations in the structure are located, on average, within the n octahedral interstitials, however significantly displaced offcenter toward the triangular face connecting to the adjacent tetrahedral vacancy, meaning that the octahedral sites are strongly involved in conduction. Ionic Conductivity of hcp-Na2B12H12‑xIx. The nonpercolating nature of T interstitials in a hcp requires the participation of O-interstitials in the conduction mechanism of 5011

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Inorganic Chemistry during the first heating cycle the conductivity reaches >0.1 S cm−1 (Figure 5, left), stabilizing at a somewhat lower value after the following heating cycles, an effect that has been observed previously.19 The determined activation energy of 140 meV is among the lowest ever reported for any Na-solid electrolyte,32 approaching liquid-like values, and certainly a new champion for the borane family. As recently stated in the context of Li solid electrolytes, the bcc lattice, where conduction proceeds exclusively via T−T motion, is regarded as a favorable environment for ionic conduction.14 The results presented therein8 were calculated on anionic sulfide lattices and determined the activation energy for hops involving the Osite in ccp or hcp lattices to be a factor 2−3 higher than those involving only T vacancies in bcc. However, the increased activation energy was owed to the energetically penalized octahedral configuration of Li in the sulfide lattice rather than the actual structural bottleneck. Computational Studies. Further insight into the factors promoting ionic mobility in the novel hcp metal borane electrolyte were obtained by a number of ab initio calculations in the framework of DFT and MD. In particular, this allowed us to investigate (i) the energy landscape of the Na-closo-borane phase diagram, (ii) the factors stabilizing the hcp in this material, and (iii) the sodium mobility in bcc, ccp, and hcp closoborane lattices. The mutual stability of the different relevant packing schemes and polymorphs was investigated by inspecting the optimized ground state energy as a function of unit cell volume. To this end, we used Na2B12H12 and Na2B12H11I chemical compositions and modeled them in different polymorphs. The volume per atom is a crucial parameter in ionic mobility. We have shown above that the experimental unit cell volume of the partially halogenated hcp material is approximately 1% larger than that of ccp Na2B12H12 at the same temperature. Since the packing density of anions is identical in both close packing schemes (74%) this observation motivates us to investigate volume effects in detail by ab initio calculations. We have systematically minimized energies of different lattice configurations based on bcc, ccp, and hcp anion lattices where Na is distributed in T- and O-sites. For each anion configuration at least four different cation distributions were considered, forming one class. The volume of these structures was fixed to the value of 280 Å3 per formula unit. In all cases the optimization of internal atomic positions was followed by simulated annealing runs (heating at 60 K/ps up to 300 K followed by cooling with at 20 K/ps, ground state structural optimization with conjugate gradient method). This procedure was repeated on volumes 262, 250, 240, 230, 220, 210, and 200 Å3 per formula unit. Additional calculations were performed for the experimental low temperature structure with P21/n symmetry. The static ground state energy for all structures is presented in Figure 6. The ground state is represented, for a volume of 220−230 Å3, by the distorted ccp lattice, which corresponds to the low temperature structure. The most stable configuration is ccp for volumes above 250 Å3 per formula unit, while it changes to bcc below 210 Å3. The hcp structure is never the preferred atomic configuration with respect to the ground state energy. On a side note it should be noted that the entropic difference between the generic hcp and ccp configurations is a mere ∼0.001kBT per sphere,33,34 in favor of the hcp lattice (lower in energy). For all structures sodium cations are located in the vicinity of tetrahedral interstitial sites: for the ccp Na+ occupy all available

Figure 6. Calculated ground state energy differences of different anion packing variants and volumes per formula unit. Data for Na2B12H12 are presented with shaded regions, and those for Na2B12H11I are shown with dashed lines.

sites; however, they are displaced from the ideal positions at the tetrahedron center. In the bcc structure the quasi-tetrahedral interstitial sites are located close to each other, two cations are located inside the coordination tetrahedron, and two are in close proximity of the tetrahedral faces. Common to all optimized hcp-based structures is that cations are located on the faces of the coordination tetrahedra. Several Na distributions within the T−T trigonal bipyramid are energetically degenerated. Such a static picture indicates the presence of partially occupied sites for Na at finite temperature. To study the influence of structure halogenation, an additional set of calculations was performed to determine, in a first instance, the stability of isolated [B 12 H 12‑x I x ] 2− considering the reaction enthalpy of B12H12 + x /2I 2 → B12H12 − xI x + x /2H 2 (x = 1, 2, 3) (1)

The energy for all possible iodine distributions within an anion was calculated, and the relative stability of the iodinated molecule is presented in Figure S20. The energy of the iodinated closo anion increases with iodine occupation, the most stable therefore is the [B12H11I]2− anion, which was considered for the calculations of the solid state halogenated Na2B12H12 structures. The relative stability of the solid state system with [B12H11I]2− and [B12H10I2]2− anions indicates an even stronger preference for single-site halogenation (see Figure S21). The ground states of Na2B12H11I were studied in an identical manner as described above for the homoleptic compound, in the volume range between 230 and 280 Å3 per formula unit. The anion substitution results in a dramatic change of the relative stability between ccp, bcc, and hcp anion packing. The hcp structure becomes the most stable in the considered volume range, with a calculated optimal volume around 250 Å3 per formula unit (dashed lines in Figure 6). In the substituted structures the large and heavy I atom is located in the vicinity of the octahedral lattice sites. The bcc anion lattice is energetically penalized due to the partial overlap of tetrahedral and octahedral sites. The ground state structures of all iodine substituted samples are shown in Figure 7d−f. The DFT optimized hcp structures do not show exactly the same cation distribution as the experimental ones. This must be understood in view of the partial occupation of Na sites. In the 5012

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Figure 7. Panels a−c show MD simulations for the three different packing (hcp, bcc, and ccp respectively) in the optimized unit cell at a volume of 262 Å3, Na atoms in light blue, boron atoms in light green. Panels d−f show equivalent static structures as guide for comparison. Hydrogen atoms are omitted for clarity, and iodine is represented by the purple sphere.

Figure 8. MD simulation for hcp structure in (a) Na2B12H12, (b) Na2B12H11I, and (c) NaCsB12H12 at a volume/f.u. of 280 Å3. Single-cation structures show a continuous 3D Na+ diffusion, whereas in the double-cation case Cs+ has the effect of limiting the Na+ long-range diffusion.

indeed cation diffusion is observed above 500 K for this structure, see Figure 7. In the hexagonal lattice cations move along the c axis up and down only between the two tetrahedra sharing a common face and forming a trigonal bipyramid, this corresponds to the above-discussed T−T hops. A second type of motion occurs in this lattice, consisting of jumps from one trigonal bipyramid to the other through an octahedral void. This motion corresponds to T−O−T hops, percolates the lattice, and shows that the hcp lattice is favorable for cation motion already at volumes smaller than those observed experimentally. The situation is even more favorable at higher experimental volumes. Both experimental approaches of stabilizing the hcp anionic sublattice (cation mixing and iodination) as well as Na2B12H12 were studied by MD with respect to the Na+ mobility at a volume of 280 Å3 per formula unit (experimental volume for hcp at 443 K). The calculations are in good agreement with experimental observations (Figure 8). By substituting [B12H12]2− with [B12H11I]2− the axis of anion rotation is locked and defined by the position of I in the octahedral site. The latter is still accessible for Na as there is an attractive interaction between negatively charged iodine and Na+. T−O−T hops creating 3D conduction pathways are observed for this structure; see Figure 8b. For NaCsB12H12 our MD simulations indicate neither sodium nor cesium mobility in the temperature range between

optimized structures only one subset of available sites is occupied, but as they are separated by rather low energy barriers (as discussed below), their occupation is equally likely. The actual mobility of sodium was investigated ab initio with MD calculations and nudged elastic band (NEB), a method for transition state searches. MD calculations were performed for all lattice types ccp, bcc, and hcp at a constant volume of 262 Å3 per formula unit and in the temperature range between 300 and 600 K. For each system an 8-fold supercell was considered, the smallest edge of this cell amounting to 10.38 Å (ccp lattice). In each system the ground state configuration was slowly heated (50 K/ps) to 300 K and equilibrated at that temperature for 8 ps. Then, heating and equilibration were repeated at T = 400, 500, and 600 K. At each temperature atomic positions and velocities were collected for structural analysis within a 15 ps time interval. Figure 7 shows the projection of the instantaneous atomic positions collected over 5 ps at 300 K on the unit cell. In all structures [B12H12]2− anions rotate on this time scale. In the ccp structure cations remain confined in the coordination tetrahedra for all simulation temperatures. Occasional jumps to the octahedral interstitials are observed, especially at higher temperatures. Cationic motion between two coordination tetrahedra via the common triangular face exists for bcc and hcp structures. In the body centered lattice this motion percolates the lattice and 5013

DOI: 10.1021/acs.inorgchem.7b00013 Inorg. Chem. 2017, 56, 5006−5016

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An activation energy of 140 meV for Na+ diffusion has been observed in hcp iodinated Na2B12H12, above 360 K, which is among the lowest ever reported among Na-solid electrolytes, approaching liquid-like values. The packing of closo-borane lattices is sensitive to disorder. We assume that this disorder will soon become a parameter that is controllable thanks to ionic substitution schemes. The synthesis of this material is achieved by low-temperature green methods without solvent and is thus sustainable and economical. The remarkable ionic conductivity reported for hcp sodium closo-borane does not require special treatment as many conventional oxide or sulfide electrolytes, and the favorable soft mechanical properties are beneficial regarding battery assembly.

300 and 700 K. In all simulations performed within this temperature range Na remains confined in the tetrahedral void and solely the T−T hops between the two face sharing tetrahedra, i.e., within a trigonal bipyramid, are observed. Cesium is confined in the octahedral void and increasing temperature the amplitude of cation motion increases to an extent that it escapes and returns back to the initial interstitial site. Thus, no jumps of Na+ between distinct trigonal bipyramids are observed, since the octahedral void occupied by Cs+ blocks this type of motion. In order to gain further understanding of the exceptionally low activation energy of 140 meV observed in the hcp halogenated phase, NEB calculations were performed for the hcp Na2B12H12 at a volume of 280 Å/f.u. Two independent sets of calculations were performed in order to determine the energy barriers for Na+ hops: (i) considering the stoichiometric composition and (ii) circumventing the Na−Na repulsion (removing one sodium and compensating for a background charge of −1e). Two classes of barriers resulted: these related to T−T hops (confined into the trigonal bipyramid) and the other related to Na hops via T−T edge or T−O face, effectively resulting in a T−O−T hop process that percolates the hcp lattice. For the stoichiometric system the T−T hop involves two cations that move simultaneously in opposite directions. This correlated motion allows the cations separation to be larger than 4.2 Å (Figure S22). Once one Na is removed, the hopping process is represented by a jump through the triangular face into the trigonal bipyramid. Both processes exhibit an energy landscape with corrugation lower than 0.15 eV (Figure S22). The second class of barriers (T−O−T hop) has a significantly higher barrier (0.4 eV) when the stoichiometric system is considered. However, once the vacant Na+ site is present, the lowest energy pathway has a barrier 95%), NaBH4 (>96%), NaI (>99.5%), and NaCl (>99.5%) were purchased from Sigma-Aldrich, CsBH4 (>95%), Na2B12H12 (>99.5%), and K2B12H12 (>99.5%) from Katchem, and Cs2B12H12 (>98%) from Strem. Anhydrous Li2B12H12 was obtained by heating Li2B12H12·0.4H20 (Katchem) in a vacuum at 513 K for 18 h. Partially Iodinated closo-Boranes Na2(B12H12‑xIx). Substituting H in B12H12 with I to synthesize the partially iodinated closo-boranes Na2(B12H12‑xIx) can be prepared using two different routes: wetchemistry using solvent and wet-milling. Route 1. The preparation method as described by Knoth et al. (1963)35 was adopted. The starting materials were combined according to the following reaction:36

Na 2B12H12 + I 2 → Na 2B12H12 − x I x + x HI Iodine (1 mmol) in 2 mL of methanol was added gradually to a solution of Na2B12H12 (1 mmol) in 0.5 mL of water +1.5 mL of methanol. The iodine color disappeared immediately. Hydrogen iodide was removed by evaporating the reaction mixture to dryness and single crystal Na2(B12H12‑xIx).2H2O was obtained. The single crystal is hydrated since it was extracted from the sample after leaving several days in air to evaporate the solvent. Efforts to grow anhydrous single crystal were not successful. Route 2. The synthesis done via water-assisted ball milling (wetmilling) leads to well-crystallized Na2(B12H12‑xIx) with excess of NaI. Reactants NaI (0.149 g., 0.001 mol) and Na2B12H12 (0.187 g., 0.001 mol) were dissolved in 1 mL of water and milled at 400 rpm for total milling time of 10 min. After milling, two-step drying was performed: first by heating at 373 K in a dynamic vacuum to remove the water, then by annealing at 411 K in an autoclave (1 atm starting pressure Argon). The resulting solid was manually ground to produce fine powder. The extent of the iodination (value of x) was unknown and might be different for each synthesis. Current efforts to obtain full structural data, including the stereochemistry of the anion, have not been successful. However, the presence of iodine isomers in the closo anion of the hydrated sample is probed by single crystal diffraction, while powder diffraction of the dried samples provides the information about anion packing which is the main objective of this work. Mixed-Cation closo-Boranes (A)(A′)B12H12 (A = Li, Na, A′ = K, Cs). The reactants were mixed in nominal composition according to Table 1 and milled at 400 rpm in a Fritsch Pulverisette 7 premium line planetary ball mill in a two-step milling process of 30 repetitions where milling intervals of 2 min are followed by breaks of 2 min to avoid overheating of the sample in the grinding bowl. The powder-to-ball mass ratio was approximately 1:50. In some cases, heat treatment such as quenching led to stabilization of very disordered phases (LiCsB12H12). Synchrotron Radiation X-ray Powder Diffraction (SR-XPD). The data used for crystal structure solution and refinements were collected at the Swiss-Norwegian Beamlines of ESRF (European

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CONCLUSION A recently published article refers to the bcc lattice as having favorable topology and energy barriers for Li-superionic conduction. Herein we report that in the family of closoboranes the octahedral site is not energetically penalized as is reported for Li-conducting sulfides. Hence, the ccp and hcp lattices, where a succession of jumps communicating between tetrahedral and octahedral sites is necessary to allow for high mobility, become very good candidates, given the seemingly stable coordination of Na[6] and Na[4]. The most favorable candidate, given the specific O−T motion, is the hcp variant. It has been shown herein that both cationic as well as anionic substitution can stabilize a hcp of [B12H12]2− anions. With these findings all the different commonly encountered packing types are reported for metal closo-boranes. The marginal differences in optimized energies indicate that meticulous crystal engineering of metal boranes can lead to a targeted stabilization of a specific packing variant. Only one of these approaches is, however, successful in terms of ionic conductivity. A second cation or a too invasive anion modification would lead to fill up the octahedral site, necessary to create a 3D conduction path involving T−O−T hops. 5014

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Inorganic Chemistry Synchrotron Radiation Facility, Grenoble, France). Data were recorded on a Dectris Pilatus M2 detector at a wavelength of 0.8170, 0.8193, 0.8198, 0.8210, 0.6888, and 0.6985 Å. The temperature was controlled with a hot air blower, and 2D images were integrated and treated with FIT2D37 or with the locally written program Bubble. High resolution data used for phase indexing were obtained at the Materials Science Beamline of the SLS (Swiss Light Source, Villigen, Switzerland) on a curved MYTHEN-II silicon strip detector at a wavelength of 0.7749 and 0.7751 Å, temperature was controlled with a hot air blower. For all measurements, the samples were sealed into borosilicate capillaries of diameter 0.5 mm (under Argon atmosphere), which were spun during data acquisition. The wavelengths were calibrated using NIST SRM660b LaB6 standard. The crystal structures presented herein were solved ab initio using the software FOX38 and refined with the Rietveld method using TOPAS.39 All the visualized structures, including MD and NEB simulation, were visualized using VESTA.40 Conduction Path Analysis. The possible conduction paths accessible to mobile Na-species were first determined by using Voronoi-Dirichlet Polyhedra (VDP) analysis implemented in the program package TOPOS. In this study, which is a geometrical approximation, [B12H12]2− and [B12H12‑xIx]2− group is represented by a single sphere, and the possible hop was constructed by connecting voids available in the close-packed lattice. The hopping channel is characterized by the radius of bottleneck (Rb), calculated as the maximal sphere mutually tangent to the three anions which define the triangular face of the interstitial site coordination polyhedra (octahedral O or tetrahedral T). The radius of the framework anion and the mobile cation is treated as Ranion = 2.9 Å and cation RNa = 1.16 Å. Electrical Conductivity Measurement. Ionic conductivities were measured using HP4192A LF impedance analyzer (frequency range 10 Hz−10 MHz, applied voltage 10 mV). Temperature-dependent impedance measurements were performed on powder samples in a symmetric cell configuration. The powder was cold-pressed into a pellet (diameter 6.35 mm, thickness ∼0.8−1.5 mm corresponds to 85−90% of the theoretical density) and spring loaded between gold electrodes placed inside Novocontrol BDS 1200 sample cell for air sensitive materials. Bulk conductivities were derived from interpreting circular arcs and/or spike intercept on the x-axis. Ab Initio Calculations. The calculations were performed within DFT formalism, and the generalized gradient approximation (GGA) for the exchange correlation functional.41 A plane-wave basis set was used and the valence configurations 1s1 for H; 2s22p1 for B; 2p63s1 for Na, and 5s25p5 for I were represented by projector-augmented wave (PAW) potentials42,43 as implemented in the Vienna Ab initio Simulation Package (VASP).44−46 The plane wave cutoff was 500 eV for all calculations, and the k-points samplings assuring total energy convergence within 1 meV per formula unit for static calculations of the ground state. The ground state electronic density was determined by iterative diagonalization of the Kohn−Sham Hamiltonian and Gaussian smearing 0.05 eV was applied. For all calculations parametrized dispersion correction was used.47 The constant volume MD studies were performed with NoseHoover thermostat,48 and the time step 0.6 fs was used for integration of the equations of motions. Single k-point was used for MD calculations. All MD calculations were performed in an 8-fold supercell (208 atoms). For simulated annealing the same parameters for integration of the equations of motion were used in the conventional unit cell. The barriers for cation hopping were calculated with climbing image NEB method49 in the same supercell as MD simulations.

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Rietveld refinement, single crystal structure and EDX data, in situ SR-PXD, structure stability, bottleneck and energy barrier calculations (PDF) CIF files (ZIP)

AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]. ORCID

Radovan Č erný: 0000-0002-9847-4372 Present Address #

EPFL ISIC-GE-VS, Rue de l’Industrie 17, Case postale 440 CH-1951 Sion. Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS This work was supported by the Swiss National Science Foundation. The authors acknowledge the Swiss-Norwegian Beamlines of ESRF and the Materials Science beamline of the SLS for the allocation of beamtime and excellent support with the data collection. Support by National Science Center Project 2015/17/B/ST3/02478, Sinergia Project CRSII2_160749/1 and CPU allocation at PL-Grid are kindly acknowledged.

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REFERENCES

(1) Coors, W. G.; Boxley, C.; Robins, M.; Eccleston, A. Low temperature molten sodium secondary cell with sodium ion conductive electrolyte membrane. WO2012061823, A3, Ceramatec, 2012. (2) Hayashi, A.; Noi, K.; Sakuda, A.; Tatsumisago, M. Superionic glass-ceramic electrolytes for room-temperature rechargeable sodium batteries. Nat. Commun. 2012, 3, 856. (3) Wang, Y.; Wang, Q.; Liu, Z.; Zhou, Z.; Li, S.; Zhu, J.; Zou, R.; Wang, Y.; Lin, J.; Zhao, Y. Structural manipulation approaches towards enhanced sodium ionic conductivity in Na-rich antiperovskites. J. Power Sources 2015, 293, 735−740. (4) 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. (5) Unemoto, A.; Chen, C. L.; Wang, Z. C.; Matsuo, M.; Ikeshoji, T.; Orimo, S. Pseudo-binary electrolyte, LiBH4−LiCl, for bulk-type allsolid-state lithium-sulfur battery. Nanotechnology 2015, 26, 254001. (6) Sveinbjornsson, 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. (7) Takahashi, K.; Hattori, K.; Yamazaki, T.; Takada, K.; Matsuo, M.; Orimo, S.; Maekawa, H.; Takamura, H. All-solid-state lithium battery with LiBH4 solid electrolyte. J. Power Sources 2013, 226, 61−64. (8) Unemoto, A.; Ikeshoji, T.; Yasaku, S.; Matsuo, M.; Stavila, V.; Udovic, T. J.; Orimo, S. Stable Interface Formation between TiS2 and LiBH4 in Bulk-Type All-Solid-State Lithium Batteries. Chem. Mater. 2015, 27, 5407−5416. (9) Paskevicius, M.; Jepsen, L. H.; Schouwink, P.; Č erný, R.; Ravnsbæk, 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. (10) Hansen, B. R. S.; Paskevicius, M.; Li, H.-W.; Akiba, E.; Jensen, T. R. Metal boranes: Progress and applications. Coord. Chem. Rev. 2016, 323, 60−70. (11) Udovic, T. J.; Matsuo, M.; Unemoto, A.; Verdal, N.; Stavila, V.; Skripov, A. V.; Rush, J. J.; Takamura, H.; Orimo, S. Sodium superionic conduction in Na2B12H12. Chem. Commun. 2014, 50, 3750−3752. (12) Udovic, T. J.; Matsuo, M.; Tang, W. S.; Wu, H.; Stavila, V.; Soloninin, A. V.; Skoryunov, R. V.; Babanova, O. A.; Skripov, A. V.;

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The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.inorgchem.7b00013. 5015

DOI: 10.1021/acs.inorgchem.7b00013 Inorg. Chem. 2017, 56, 5006−5016

Article

Inorganic Chemistry

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DOI: 10.1021/acs.inorgchem.7b00013 Inorg. Chem. 2017, 56, 5006−5016