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Fast Lithium-ion Conduction in Atom-Deficient Closo-Type Complex Hydride Solid Electrolytes Sangryun Kim, Naoki Toyama, Hiroyuki Oguchi, Toyoto Sato, Shigeyuki Takagi, Tamio Ikeshoji, and Shin-ichi Orimo Chem. Mater., Just Accepted Manuscript • DOI: 10.1021/acs.chemmater.7b03986 • Publication Date (Web): 20 Dec 2017 Downloaded from http://pubs.acs.org on December 20, 2017

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Chemistry of Materials

Fast Lithium-ion Conduction in Atom-Deficient Closo-Type Complex Hydride Solid Electrolytes Sangryun Kim,*,† Naoki Toyama,† Hiroyuki Oguchi,‡ Toyoto Sato,† Shigeyuki Takagi,† Tamio Ikeshoji,† and Shin-ichi Orimo,†,‡ †

Institute for Materials Research, Tohoku University, Katahira 2-1-1, Aoba-ku, Sendai 980-8577, Japan WPI-Advanced Institute for Materials Research (WPI-AIMR), Tohoku University, Katahira 2-1-1, Aoba-ku, Sendai 9808577, Japan

‡

ABSTRACT: Closo-type complex hydrides contain large cage-type complex polyanions in their crystal structures and thus can exhibit superior ion-conducting properties (e.g., Li and Na). However, the unique structures of complex polyanions have made it challenging to modify crystal structures, making systematic control of ion conductivity difficult. Here, we report an atom deficiency approach to enhance lithium-ion conductivity of complex hydrides. We find that lithium and hydrogen could be simultaneously extracted from Li2B12H12 by applying a small external energy, enabling the formation of atom deficiencies. These atom deficiencies lead to an increase in carrier concentration, improving lithium-ion conductivity by three orders of magnitude compared to that of a pristine material. An all-solid-state TiS2/Li battery employing atom-deficient Li2B12H12 as a solid electrolyte exhibits superior battery performance during repeated discharge–charge cycles. The current study suggests that the atom deficiency can be a useful strategy to develop high ion-conducting complex hydride solid electrolytes.

1. INTRODUCTION Recent intensive research effort has led to considerable progress in all-solid-state batteries, in which both the electrolyte and electrodes are in their solid states. The advantages of allsolid-state batteries lie in their ability to overcome the intrinsic drawbacks of conventional liquid-based batteries, such as electrolyte leakage, flammability, and limited voltage window, as solid-state electrolytes are usually non-explosive, non-volatile, and electrochemically stable, even up to ~5.0 V vs. Li+/Li.1,2 In particular, increasing safety concerns regarding the application of lithium batteries with liquid electrolytes in electric vehicles have stimulated research into allsolid-state batteries. Lithium-ion-conducting solid electrolytes are a key component of all-solid-state batteries because the stability and ionic conductivity of the solid electrolyte determine battery performance.3-8 Among the variety of materials reported to date, complex hydrides have recently attracted particular attention as a new class of solid electrolytes owing to their superior stability against lithium metal, which results from their high reducing ability, as well as their high ionic conductivity.9-11 Complex hydrides are generally denoted by M(M’xHy), where M is a metal cation and M’xHy is a complex anion. Initial interest in complex hydride ionic conductors has mainly focused on lithium borohydride (LiBH4) and related derivative materials.9 A series of closo-boranes containing complex polyanions such as [B12H12]2-, [CB11H12]-, and [CB9H10]- have recently been reported to exhibit high lithium-ion conductivities.12-14 Despite their impressive ionic conductivities, the unique complex anion structure formed by strong covalent bonding between the boron and hydrogen atoms greatly complicate

efforts to modify crystal structures, which are directly related to their ionic conductivities. For example, systematic cationic and/or anionic substitutions, which have been widely adopted to control the crystal structures (and thus ionic conductivities) of various ionic conductors,15,16 are largely inappropriate because the structural stabilities of complex anions are significantly lowered by such substitutions, leading to structural collapse. For these reasons, substitution of the complex anions themselves to form so-called mixed complex anions,9,17,18 has been explored, but usable candidates are very limited. Therefore, a more general approach is required to broaden the structure and conduction property scopes of complex hydrides. Herein, in an effort to overcome the chronic drawbacks of complex hydrides and enhance their ionic conductivities, we report the unconventional approach of introducing atom deficiencies to lithium dodecahydro-closo-dodecaborate (Li2B12H12). Our experimental results reveal that simple ballmilling generates both lithium and hydrogen deficiencies, and that the increased carrier concentration by these deficiencies increases lithium-ion conductivity significantly. Furthermore, an all-solid-battery using atom-deficient Li2B12H12 as the solid electrolyte exhibits good cycling stability during repeated discharge-charge cycles. The present investigation demonstrates that the exploitation of atom deficiencies is a viable approach to improving the ionic conductivities of closo-type complex hydride solid electrolytes. The effects of mechanical ball-milling have been previously investigated for the formation of a disordered hightemperature (HT) phase, in which fast ionic conduction is enabled predominantly by the dynamics of the complex anions.14,19 Our careful characterizations reveal the distinctly

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different ionic conduction mechanism of atom-deficient Li2B12H12 compared to those present in previous reports.

2. EXPERIMENTAL SECTION Sample synthesis. Pristine Li2B12H12: The starting material, hydrated Li2B12H12·4H2O, was purchased from Katchem. Li2B12H12·4H2O was first ground using a mortar and pestle for 15 min. To obtain anhydrous samples, the powders were subsequently dried under vacuum (< 5 × 10−4 Pa) at 498 K for 20 h. Mechanically ball-milled Li2B12H12: To synthesize mechanically ball-milled Li2B12H12, the anhydrous Li2B12H12 powder was milled using a planetary micro mill (Pulverisette 7, Fritsch) at 400 rpm for 5 h in an Ar atmosphere. The resulting powders were finally heat-treated at 423 K for 12 h under vacuum to remove residual water. Characterization. Crystal structures were characterized using X-ray diffraction (XRD, X’PERT Pro, PANalytical) by scanning in the 2θ range of 10−50° with a scan step of 0.03° and an acquisition time of 3 s for each step. The powder for the XRD measurements was loaded into a thin-walled glass capillary under Ar atmosphere and sealed with vacuum grease. Morphologies and particle sizes were analyzed using a scanning electron microscope (SEM, JSM-6009, JEOL). The lithium and hydrogen contents of the compounds were obtained using inductively coupled plasma (ICP, IRIS Advantage DUO, Thermo Fisher Scientific) measurement and a hydrogen analyzer (EMGA-621W, Horiba). The bonding states were characterized by Raman spectroscopy (Nicolet Almega-HD, Thermo Scientific). Differential thermal analysis (DTA) was performed using a Rigaku Thermo plus TG8120 system from 25 to 400 °C at a ramping rate of 5 °C min−1 under an Ar flow. Ionic conductivities were measured using Li-symmetric electrodes by the AC impedance method over a temperature range of 298−353 K with applied frequencies in the range of 4−1 MHz using a frequency response analyzer (3532-80, Hioki). An electronic conductivity was measured via the DC method using an electrochemical measurement system (1470E, Solartron Analytical). The electrochemical stability was evaluated by cyclic voltammetry (CV, 1470E, Solartron Analytical) using stainless steel (SS)/ball-milled Li2B12H12/Li cell at a scan rate of 5 mV s−1 and a scan range of −0.1 to 5 V (vs. Li+/Li). Battery test. A composite electrode consisting of commercial TiS2 (Sigma-Aldrich) and ball-milled Li2B12H12 as a cathode material was fabricated. The TiS2 and ball-milled Li2B12H12 powders were weighed at a 2:3 mass ratio and mixed using an agate mortar and pestle. For cell fabrication, 30 mg of the ball-milled Li2B12H12 powder was first placed into an 8-mm-diameter die and uniaxially pressed at 60 MPa. Subsequently, 5 mg of the composite positive electrode powder was transferred onto the pressed electrolyte still present in the die and uniaxially pressed at 240 MPa to obtain one single pellet comprising of the cathode and solid electrolyte. Lithium metal (Honjo Metal) was used as a negative electrode and placed on the opposite side of the positive electrode. The assembled all-solid-state TiS2/ball-milled Li2B12H12/Li metal cell was placed in a stainless steel electrochemical cell with an 8-mm-diameter Teflon guide, as schematically illustrated elsewhere.20 These cell components were assembled in an argon-filled glovebox. Cycling tests

were carried out in the voltage range 1.6−2.7 V (vs. Li+/Li) with an applied current of 0.05C (1C = 239 mA g−1) at 80 °C using a battery tester (580 Battery Test System, Scribner Associates). First-Principles Molecular Dynamics Simulation. The total energy changes due to atom deficiencies were estimated by first-principles molecular dynamics (FPMD) simulation at 300 K. In the FPMD simulations, an electronic structure calculation code implemented in the Vienna ab initio simulation package (VASP) using density functional theory (DFT) was used with plane-wave basis sets and ProjectorAugmented-Wave (PAW) pseudo potentials under periodic boundary conditions. A Perdew-Burke-Ernzerhof (PBE) functional21 was used for the exchange correlation with a generalized gradient approximation. The cutoff energy was 320 eV for wave functions, and k-point sampling was 1 × 1 × 1. The starting structure had 32 Li2B12H12 formula units (f.u.) (Z = 4, 2 × 2 × 2 supercell), and the formation energies were calculated from the time average total energy E(LiaBbHc) using the following equations: ELi = E(Li63B384H384) + E(Li) − E(Li64B384H384) EH = E(Li64B384H383) + E(H2)/2 − E(Li64B384H384) ELiH = E(Li63B384H383) + E(LiH) − E(Li64B384H384) where E(Li) and E(LiH) are the ground-state energies of Li metal (BCC) and LiH (rock-salt), respectively. E(H2) is the ground-state energy of a H2 molecule in a vacuum. The total simulation time was 12 or 16 ps and the last 6-ps data after the steady state was almost achieved were used to calculate each average total energy.

3. RESULTS AND DISCUSSION Figure 1 shows the crystal structure of Li2B12H12 consisting of Li+ cations and [B12H12]2- complex polyanions. In typical alkali cation dodecahydro-closo-dodecaborates (M’’2B12H12, M’’ = alkali cation), the M’’+ cation is tetrahedrally surrounded by four [B12H12]2- anions. For Li2B12H12, the Li+ cation is displaced 1.82 Å from the center of the tetrahedral site.22 It is well known that large and quasi-spherical [B12H12]2- polyanions provide facile ionic conduction channels with low activation energy barriers for diffusion.14 Atom-deficient Li2B12H12 was prepared via a simple mechanical method using high-energy ball-milling, which has been widely used to introduce atom deficiencies into various battery materials.23-25 Structural changes upon ball-milling

Figure 1. The crystal structure of Li2B12H12 consisting of Li+ cations and [B12H12]2- complex polyanions in ball-and-stick (left) and polyhedral (right) representations. Red, green, and blue spheres represent Li, B, and H, respectively.


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Table 1. Lithium and hydrogen contents per chemical formula for the pristine and ball-milled samples.

Figure 2. (a) XRD patterns of pristine and ball-milled Li2B12H12 and (b and c) their SEM images.

were first investigated using XRD measurements (Figure 2a). The diffraction peaks of both the pristine and ball-milled samples are well indexed based on a cubic unit cell with the space group Pa3 (No. 205).22 As neither impurity phases nor peak splits are detected for the ball-milled Li2B12H12, the peak shifts to lower angles observed indicates an expansion of the lattice without perturbing the crystal symmetry. The lattice parameters of the pristine and ball-milled samples were calculated to be a = 9.5853(4) and a = 9.6283(8) Å, respectively. SEM analyses indicate no significant differences in morphology or particle size between the samples (Figures. 2b, 2c, and S1). The lattice change upon ball-milling (Figure 2a) may be caused by the atom deficiencies. As direct evidence, the lithium and hydrogen contents, as summarized in Table 1, were investigated using ICP analysis and a hydrogen analyzer. Interestingly, the compositional characterizations revealed decreased contents of both lithium and hydrogen in the ballmilled sample, indicating that atom deficiencies are formed upon ball-milling. These atom deficiencies are expected to cause substantial lattice evolution that gives rise to changes in atomic ordering and lattice distortion, both of which are preferable to stabilize the crystal structure. The decrease in the intensity of the (021) peak after ball-milling (Table S1) verifies the rearrangement of lithium atoms at multiplesites.26 In addition, the local distortion of B-H bonding in the [B12H12]2- polyanions was confirmed by Raman characterization (Figure S2). To thermodynamically evaluate the atom deficiencies from Li2B12H12, the formation energies of atom-deficient phases were estimated using first-principles molecular dynamics (FPMD) simulations. The starting structure had 32 Li2B12H12 formula units (f.u.), and three types of formation

Li content

H content

Theoretical

2

12

Pristine

2.01(1)

12.03(5)

Ball-milled

1.92(1)

11.69(14)

energies were calculated by introducing different atom deficiencies, i.e., 1Li (ELi: Li-deficient), 1H (EH: H-deficient), and 1Li + 1H (ELiH: Li- and H-deficient) to the starting structure. The ELi, EH, and ELiH values calculated per f.u. were 2.9 (= 0.03 eV), 4.2, and 2.2 kJ mol−1, respectively. The error estimated from the standard deviation of moving average with a 1-ps window was 1.5 kJ mol-1. These results indicate that small amounts of lithium and hydrogen could be extracted as various types of compounds such as Li, H2, and LiH by applying a small external energy, enabling the formation of atom deficiencies. Our experimental and theoretical results, which revealed atom deficiencies of both lithium and hydrogen, are distinct from previous computational expectations, in which only the cation (Li or Na) deficiency was used.27,28 Atom deficiencies in complex hydrides are very rare, and have been reported for only a few closo-boranes.26,29 The important feature of closo-boranes that differentiates them from LiBH4-related compounds is that the complex polyanions comprise multiple boron atoms. In such closo-type complex polyanions, the boron and hydrogen atoms form a particular covalent bonding pattern (boron-boron and boronhydrogen), leading to the robust cage-type polyanionic structure.30 Thus, the stable structure of the complex polyanions may contribute to the structural stability of the H-deficient phase. In addition, the similarity of ELi, EH and ELiH suggests that a certain amount of lithium deficiency would compensate for some amount of H deficiency. Analyses to identify detailed mechanisms for the formation of these atom deficiencies are underway. The ionic conductivity of ball-milled Li2B12H12 was assessed using impedance measurements. In order to investigate the role of ball-milling, the pristine sample was also tested. Figure 3a shows Arrhenius plots of the ionic conductivities measured in the temperature range 298−353 K. The impedance profile presents one clear semicircle, which can be interpreted as the parallel combination of a resistance and a capacitance (Figure S3). Empirically, capacitances of the bulk and grain boundary for ionic conduction are in the order of 10−12 and 10−9 F, respectively.31 The capacitance at 30 °C (= 303 K) was calculated to be 6.4 × 10−11 F, hence the observed impedance corresponds to a total resistance (bulk + grain boundary). The total conductivity (σ) at 30 °C and activation energy (Ea) of pristine Li2B12H12 are 2.5 × 10−8 S cm−1 and 47.4 kJ mol−1, respectively. These results are in good agreement with those of a previous report.14 Importantly, it appears that the atom deficiencies in Li2B12H12 provide quantitatively independent changes in the ionic conductivity and activation energy. While ball-milled Li2B12H12 exhibits an ionic conductivity (2.0 × 10−5 S cm−1 at 30 °C) three orders of magnitude higher than that of the pristine counterpart, its activation energy (44.1 kJ mol−1) is



Figure 3. (a) Arrhenius plots of the conductivities of pristine and ball-milled Li2B12H12. (b) Discharge-charge profiles of an all-solid-state battery consisting of a TiS2 cathode, a ball-milled Li2B12H12 solid electrolyte, and a Li metal anode at 80 °C and 0.05C.

almost the same. These behaviors can be understood from the Arrhenius equation: σ = σ0 exp (-Ea/RT), where σ (S cm−1) is the conductivity, σ0 is the pre-exponential parameter (S cm−1), Ea is the activation energy (kJ mol−1), R is the universal gas constant (kJ mol−1 K−1), and T is the absolute temperature (K). The pre-exponential parameter is proportional to the net concentration of carriers (lithium and vacancy).32 These relationships indicate that the improved ionic conductivity after ball-milling is due to an increase in carrier concentration. The electronic conductivity of the ball-milled Li2B12H12 was on the order of 10−9 S cm−1; hence the conductivities estimated from the impedance measurements are attributed to lithium ion conduction. On the basis of our structural characterizations, it is speculated that the increased carrier concentration originates from the atom deficiencies. The lithium deficiencies can allow the formation of additional lithium sites as well as a change in lithium arrangement, both of which correlate with the carrier concentration. The lowered intensity of the (021) peak after ball-milling, which is confirmed from the XRD analyses (Figure 2a and Table S1), reflects the rearrangement of the carriers (lithium and vacancy) at multiple sites, increasing the carrier concentration. A previous structural study reported that the reduced intensity of the (021) peak indicates lithium arrangement at multiple sites (8c and 24d) with the lowered occupancies.26 Other lithium sites can also be generated by the hydrogen-deficient and/or distorted complex anions. Similar conductivity changes involved with the carrier con-

centration have been observed and well explained for perovskite-type lithium ionic conductors.15,32,33 The ball-milled Li2B12H12 was examined as a solid electrolyte in an all-solid-state battery. The electrochemical stability was first evaluated from the CV of an SS/ball-milled Li2B12H12/Au cell at a scan rate of 5 mV s−1 and a scan range of −0.1 to 5 V (Figure S4). The cell displayed reversible cathodic and anodic currents around 0 V, which correspond to lithium deposition and dissolution, respectively. In addition, no oxidation currents were detected in the scanned voltage range, thus indicating the wide electrochemical stability of the ball-milled Li2B12H12. Figure 3b shows the dischargecharge curves of the all-solid-state battery, fabricated using a TiS2 cathode, a ball-milled Li2B12H12 solid electrolyte, and a lithium metal anode. When galvanostatically cycled in the voltage range of 1.6−2.7 V (vs. Li+/Li) at 80 °C and 0.05C (1C = 239 mA g−1), the all solid-state battery exhibited a first discharge capacity of 228 mAh g−1 and good capacity retention over 20 cycles, demonstrating that ball-milled Li2B12H12 can be used as the solid electrolyte in all-solid-state batteries. However, the cell that utilized a pristine Li2B12H12 as the solid electrolyte delivered smaller capacity and worse capacity retention over repeated discharge−charge cycles.34 It has been reported that a disordered HT phase of closoboranes can be produced by ball-milling.14 In order to clarify the relationship between the ball-milled phase and the HT phase, their structures were compared. It is known that Li2B12H12 undergoes transition to the HT phase at ~355 °C, which is accompanied by a significant rise in ionic conductivity.14,26 During this phase transition, the lattice parameter is significantly expanded (8.7%) while maintaining the crystal symmetry. Comparison of the HT phase (extrapolated lattice parameter at room temperature a > 9.9 Å) and the ball-milled phase (a = 9.6283(8) Å) reveals a large difference in lattice parameters, implying that both phases represent different crystal structures. More precise differences can be verified by focusing on the (021) peak, the intensity of which is related to the lithium arrangement. In the case of the HT phase, a reduced (or invisible) (021) peak indicates a completely disordered lithium arrangement after phase transition (Table S1).26 In contrast, the ball-milled Li2B12H12 presents a preserved (021) peak, reconfirming the structural disagreement between the phases. More critically, a transition to the HT phase is detected for the ball-milled Li2B12H12. The DTA data for the ball-milled Li2B12H12 present clear endothermic and exothermic peaks at ~310 and 275 °C upon heating-cooling cycling (Figure 4), which originate from the reversible transitions to and from the HT phase. These results, once again, indicate that the ball-milled Li2B12H12 is distinct from the HT phase. The slightly lower transition temperature for the ball-milled Li2B12H12 compared to that of the pristine material is presumably due to the effects of the increased entropy change. The fast ionic conduction in the atom-deficient complex hydrides observed in the present study is of importance because it can be utilized as a design principle for developing closo-type complex hydride solid electrolytes. Wellcontrolled atom deficiencies can further improve ionic con ductivities, as previously demonstrated in various lithium ionic conductors.8,15,35 This approach can be applied to a variety of hydrides with cage-type complex polyanions. To achieve this, systematic routes to forming atom deficiencies


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ACKNOWLEDGMENT The authors would like to thank Mr K. Sato, Ms H. Ohmiya, and Ms N. Warifune for technical assistance, and the use of SR16000 supercomputing resources at the Center for Computational Materials Science of the Institute for Materials Research, Tohoku University. This work was supported by JSPS KAKENHI (Grant Numbers 17H06519, 16K06766, 17K19168, 17K18972, and 25220911), Collaborative Research Center on Energy Materials in IMR (E-IMR), and Target Project 4 of WPI–AIMR, Tohoku University.

REFERENCES

Figure 4. DTA curves of pristine and ball-milled Li2B12H12 during a heating-cooling cycle.

need to be established. In addition, investigation of atomdeficient complex hydrides can be useful for understanding the ionic conduction in related materials of the same category. Since the HT phases of closo-type complex hydrides exhibit extremely high ionic conductivities, a variety of approaches, such as mechanical ball-milling,14 the use of mixed complex anions,36 and atomic substitution in complex anions,13 have been attempted to stabilize the HT phase at lower temperatures. However, after the given treatments, the materials showed large disagreements in conductivities as compared to the pristine materials as well as non-linear Arrhenius profiles, implying the possibility of other factors influencing the ionic conductivities. In this regard, the atom deficiencies might be a viable approach to revealing the unexplained conduction mechanisms of closo-type complex hydrides.

4. CONCLUSIONS In summary, this investigation introduces a new strategy to enhance the ionic conductivity of closo-boranes, i.e., the formation of atom deficiencies. The atom deficiencies in Li2B12H12 result in high carrier concentration originating from multiple lithium sites and increased vacancies, thus improving lithium-ion conductivity. Highly important is the fact that the proposed high ionic conductivity of the atomdeficient phase can be utilized together with the effects of the HT phase. Thus, the introduction of atom deficiencies provides useful insight into strategies that can be applied to developing complex hydride solid electrolytes.

ASSOCIATED CONTENT Supporting Information The Supporting Information is available free of charge on the ACS Publications website. Additional characterization data (SEM, ratios of XRD peaks, Raman, impedance, and CV) (PDF)

AUTHOR INFORMATION Corresponding Author *E-mail: [email protected]

Notes The authors declare no competing financial interest.

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