Enhanced Li Ion Conductivity in LiBH4−Al2O3 Mixture via InterfaceEngineeringYong Seok Choi,†,‡ Young-Su Lee,*,† Dong-Jun Choi,† Keun Hwa Chae,§ Kyu Hwan Oh,‡

and Young Whan Cho†

†High Temperature Energy Materials Research Center, Korea Institute of Science and Technology, Seoul 02792, Republic of Korea‡Department of Materials Science and Engineering, Seoul National University, Seoul 08826, Republic of Korea§Advanced Analysis Center, Korea Institute of Science and Technology, Seoul 02792, Republic of Korea

*S Supporting Information

ABSTRACT: A new solid-state Li ion conductor composed of LiBH4 andAl2O3 was synthesized by a simple ball-milling process. The elementdistribution map obtained by transmission electron microscopy demonstratesthat the LiBH4 and Al2O3 are well mixed and form a large interface after ball-milling. The ionic conductivity of the mixture reaches as high as 2 × 10−4 Scm−1 at room temperature when the volume fraction of Al2O3 isapproximately 44%. The ionic conductivity of the interface between LiBH4and Al2O3 was extracted by using a continuum percolation model, whichturns out to be about 10−3 S cm−1 at room temperature, being 105 timeshigher than that of pure LiBH4. This remarkable rise in conductivity isaccompanied by the lowered activation energy for the Li ion conduction inthe mixture, indicating that the interface layer facilitates Li ion conduction.Near-edge X-ray absorption fine structure analysis reveals the presence of B−O bondings in the mixture, which was not detected by X-ray diffraction. Thisdisruption of the chemical bondings at the interface may allow an increase in carrier concentration and/or mobility therebyresulting in the pronounced enhancement in conductivity. This result provides a guideline for designing fast Li ion conductorthrough interface engineering.

1. INTRODUCTION

All solid-state lithium ion batteries with solid electrolytes havedrawn much attention as an alternative to batteries usingorganic liquid electrolytes, due to their high power density andgreat thermal/mechanical stability.1−4 However, it still remainsa challenge to design suitable solid electrolytes with sufficientchemical and electrochemical stability, and also with highlithium ion conductivity. Among the various solid electrolytesdiscovered thus far,5−16 lithium borohydride (LiBH4) isconsidered one of the most promising candidates. LiBH4

possesses excellent properties including high electrochemicalstability of up to 5 V (Li+/Li), high temperature durability, andnegligible electronic conductivity.17,18 Matsuo et al.19 discov-ered that LiBH4 displays high Li ionic conductivity of ∼10−3 Scm−1 above 110 °C, which is the phase transition temperaturefrom low temperature phase (orthorhombic structure, Pnmaspace group) to high temperature phase (hexagonal, P63mc).Several researchers have fabricated all solid-state batteries withLiBH4 as a solid electrolyte and successfully demonstrated thecharge and discharge of the battery with active electrodes suchas S, LiCoO2, and TiS2 at high temperatures above 110°C.20−24

However, the ionic conductivity of LiBH4 drasticallydecreases by several orders of magnitude after phase trans-

formation to the orthorhombic phase and reduces to ∼10−8 Scm−1 at room temperature (RT). Since the high transitiontemperature is an unavoidable obstacle to adopting LiBH4 incommercial batteries, many researchers have made an effort toachieve high ionic conductivity at low temperature throughtailoring the chemical structure by doping halides,25−27

hydration,28 and combining complex anions.29−31 As ananother route to increase conductivity, our previous studydemonstrated that ionic conductivity of LiBH4 at RT can beenhanced up to 4 orders of magnitude higher by mixing withSiO2 nanoparticles.

32 Liang and colleagues first discovered sucha phenomenon in the LiI−Al2O3 composite, where the additionof Al2O3 increased the ionic conductivity of LiI by 2 orders ofmagnitude.33 These kinds of mixtures, i.e., ionic conductor withinsulating fine particles, have been termed dispersed ionicconductors and are characterized by the formation of defectiveand highly conducting space-charge layers at the interfacebetween the conductor and insulator. Similar enhancement inionic conductivity has been found when insulating matters,such as SiO2, Al2O3, and B2O3, are mixed with Li ion

Received: September 6, 2017Revised: October 29, 2017Published: November 10, 2017

Article

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© 2017 American Chemical Society 26209 DOI: 10.1021/acs.jpcc.7b08862J. Phys. Chem. C 2017, 121, 26209−26215

Cite This: J. Phys. Chem. C 2017, 121, 26209-26215

conductors, highlighting the importance of interface engineer-ing.32,34−36 Most recently, Blanchard et al.37 suggested thatnanoconfinement of LiBH4 into a mesoporous scaffold (MCM-41) can maximize the interface area between two mediums, andthe Li transport is much faster compared to bulk LiBH4. Suchinterface control is worthwhile because it can be applicable toany kinds of ionic conductors, but the main mechanism of fastLi transport at the interface has not yet been elucidated. Inorder to develop fast ionic conductors in the future, it is crucialto identify the key factors that govern Li mobility at theinterface.Herein we characterize the interface of a composite Li ion

conductor, LiBH4−Al2O3, and attempt to identify themechanism of conductivity enhancement. In order to improvethe ionic conductivity through interface engineering, weselected the γ-Al2O3 nanoparticle (primary particle diameterof 5 nm) as an additive insulator that is expected to form a largeinterface area with LiBH4. It is found that the ionic conductivityof the LiBH4−Al2O3 mixture can be raised to 2 × 10−4 S cm−1

at RT by optimizing the volume fraction of Al2O3. Comparisonwith the interface of the LiBH4−SiO2 mixture reveals superiorion conduction at the LiBH4−Al2O3 interface.

2. EXPERIMENTAL METHODS

Commercial LiBH4 powder (purity 95%, Acros) and γ-Al2O3(purity 99.99%, US research nanomaterials) were used asstarting materials. In order to reduce the particle size,commercial LiBH4 powder was ball-milled using a planetarymill (Fritsch P7) operated at 600 rpm for 2 h. As received γ-Al2O3 was dried under vacuum at 350 °C for 7 h to remove theadsorbed water and oxygen. Then, premilled LiBH4−Al2O3(LA) mixtures with five different volume fractions (19, 33, 44,54, and 73 vol % of Al2O3; see Supporting Information fordetails) were ball-milled for uniform mixing. A ball-millingprocess was conducted using the planetary mill operated at 200rpm for 30 min. All handling of the samples was performed in aglovebox filled with argon (p(O2), p(H2O) < 1 ppm).The X-ray diffraction (XRD) patterns of LA mixtures were

measured using a Bruker D8 Advance diffractometer with CuKα radiation operating at 40 kV and 40 mA at RT. Theexposure time and step size were 1 s step−1 and 0.03°,respectively. To prevent air exposure, samples were sealed withpolyimide thin-film (7.5 μm) tape in the glovebox. A highresolution transmission electron microscope (TEM, FEI Titan)operating at 200 keV was used to observe the microstructure ofthe solid electrolyte. A TEM sample with a thickness of ∼80nm was prepared using a focused ion beam (FIB, FEI Quanta3D FEG). All sample transfer processes were conducted using amobile air-lock holder. Digital Micrograph software was usedfor analysis of selected area electron diffraction (SAED) andelectron energy loss spectroscopy (EELS).The ionic conductivity of the LA mixtures was measured by

AC impedance spectroscopy using a Solartron impedanceanalyzer (SI 1260). A detailed description of the impedancemeasurement can be found in the Supporting Information. Todetect a change in bonding structure after ball-milling, nearedge X-ray absorption fine structure (NEXAFS) measurementswere conducted at 10D XAS KIST beamline in Pohang LightSource (PLS), Korea. B K-edge NEXAFS spectra werecollected in both total-electron yield (TEY) mode andfluorescence yield (FY) mode and are normalized with respectto the incident photon flux.

3. RESULTS AND DISCUSSIONFigure 1 shows the XRD patterns of LA mixtures before andafter the ball-milling process. All LA mixtures before ball-

milling (Figure 1a) display the features of starting materials,including strong peaks of LiBH4 and broad peaks of γ-Al2O3.The peak intensity of LiBH4 is proportional to the amount ofLiBH4. The broad peak around 18° comes from the polyimidethin film used to prevent air exposure. After ball-milling, theXRD patterns of LA mixtures (Figure 1b) remain almostidentical except for broadening of the LiBH4 peaks. Thebroadened peaks indicate the reduction in the crystallite size ofLiBH4, but any chemical reaction between LiBH4 and Al2O3during ball-milling is not evident in the XRD patterns.The formation of a large interface between conductor and

insulator and its percolation throughout the conductor areimportant for improving conductivity by interface engineering.Thus, we attempt to visualize the phase distribution in themicrostructure of LA 44 vol % mixture after ball-milling usingTEM. Figure 2a shows the cross-sectional image of the LA 44vol % sample. It can be seen that tiny particles exhibiting darkor bright contrast are well mixed at the nanometer scale. EELSelement mapping (Figure 2b) clarifies that Li and B signalsmainly come from the dark contrast region, while Al signals arestronger at the bright contrast area. Ring patterns in SAED(Figure 2a) match well with the calculated positions of LiBH4and Pt (deposited for surface protection during TEMsampling), but the pattern for Al2O3 is barely visible becauseit is nanocrystalline. This result is in good agreement with theXRD analysis. From this microstructure analysis, we can predictthat a large area of interface between LiBH4 and Al2O3 iscreated by ball-milling, and that would possibly provide a fastconduction path for Li ions.Figure 3 is a plot of the ionic conductivities of LA mixtures as

a function of temperature, measured by electrochemicalimpedance spectroscopy. Interestingly, even a small amountof Al2O3 (∼19 vol %) greatly improves ionic conductivity up to4 × 10−5 S cm−1 at RT, which is almost 3 orders of magnitude

Figure 1. X-ray diffraction patterns of LA mixtures (vol % of γ-Al2O3in the legend) (a) before and (b) after ball-milling process.

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higher than that of pure LiBH4. The conductivity increasescontinuously as the amount of Al2O3 increases and hits amaximum of ∼2 × 10−4 S cm−1 at RT when the volumefraction of Al2O3 reaches 44%. Notably, this is the highest valuecompared to other studies32,37 that report conductivityenhancement through interface engineering (see Figure S3for comparison). In addition, an LA 44 vol % mixture displays avery tiny change in conductivity at the phase transitiontemperature, at which the conductivity of bulk LiBH4 suddenlyrises by 3 orders of magnitude. This implies that the Li ionconduction is dominated by the interface layer with a minorcontribution by bulk LiBH4. After hitting the maximum point,the ionic conductivity decreases as the volume fraction of Al2O3increases because the large amount of insulating materialsdisrupts the percolation of the interface layer. Based onprevious literature and our own studies, we conclude that the

ionic conductivity enhancement by mixing oxide compounds(e.g., SiO2, Al2O3, and B2O3) is a universal phenomenonapplicable to a wide range of conductors.32,35−38 However, thedegree of enhancement varies in the range of 100 to 10000times higher depending on the chemical composition and sizeof the insulating materials being mixed. In our LA mixtures, thedegree of enhancement appears among the highest, and thissuggests that the interface engineering is a very effective routeto increase the conductivity of LiBH4.The conductivity data in Figure 3 can be analyzed in a

different fashion, i.e., conductivity variation as a function ofvolume fraction of Al2O3 at a specific temperature, which ispresented in Figure 4a. On top of the experimentally measuredconductivity, we have simulated the conductivity of a composite(σ(p)) as a function of the volume fraction of an insulator (p)using a continuum percolation model developed by Roman etal.39 Using their formalism, the ratio τ (= σA /σB) between theinterface conductivity (σA) and the bulk conductivity of LiBH4

(σB) can be estimated (see Supporting Information for thedetails). The simulated curves (solid lines) with τ values of 1.5× 105 (at 25 °C), 5.6 × 104 (at 62 °C), 2.3 × 104 (at 99 °C),and 37 (at 124 °C) reasonably reproduce the trend of themeasured values (solid symbols) shown in Figure 4a.We compare the interface conductivity of LA mixtures with

that of LiBH4-fumed silica (LF) mixtures in our previousstudy32 and present the result in Figure 4b. Overall, theinterface conductivity is much higher than that of bulk, being∼10−3 S cm−1 at RT, as can be immediately known from theextremely high τ value. Comparison between LA and LFmixtures reveals that the interface conductivity of LA mixturesis almost twice as high as that of LF mixtures irrespective of thetemperature. Therefore, the higher τ value (or interfaceconductivity) in the LA mixture mainly, but not fully, accountsfor the higher total conductivity of LA (see Figure S3) and therest can be attributed to the increased proportion of the volume

Figure 2. (a) TEM image of the microstructure of LA 44 vol % mixture and corresponding SAED pattern. (b) EELS element mapping of red squarearea in (a). Red, green, and blue color represent Li, B, and Al, respectively.

Figure 3. Arrhenius plots of the ionic conductivities of LA mixtures indifferent mixing ratios (vol % of Al2O3 in the legend). The gray dotsare data of pure LiBH4. All the data were obtained from the secondcooling run.

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fraction of interface due to the smaller particle size of γ-Al2O3(∼5 nm) compared to fumed silica (∼7 nm).32,36

The difference in τ values suggests that the interfaceconductivity can be varied depending on the chemicalcomposition of the oxide particles, and the LiBH4−Al2O3interface might be more efficient in promoting the diffusionof the Li ion than the LiBH4−SiO2 interface. Suwarno et al.observed a similar phenomenon that the chemical nature of thescaffolds (carbon and silica scaffolds in their study) can affectthe hydrogen mobility and interface layer thickness by changinginterfacial energy.40 In their study, the interfacial thickness ofLiBH4 with SiO2, deduced from the decreased enthalpy changeof the phase transition, is found to be larger than that withcarbon. In our analysis, the interfacial thickness was arbitrarilyfixed to 1 nm for both LA and LF mixtures, and therefore therelatively high τ value of LA mixtures may also reflect theincreased interfacial thickness in LA mixtures. Further study isnow underway to understand how the chemical composition ofan insulator alters the interface conductivity and to find outwhether other nonoxide compounds can deliver the con-ductivity enhancement to a similar degree.Since the ion conduction in solid takes place by jumping of

the ions to the adjacent vacancies or interstitial sites, thestructural evolution at the interface strongly influences theactivation energy required for Li migration. The activationenergy (Ea

σ) of mixtures was determined from the slope ofArrhenius plots of the ionic conductivities below the phasetransition temperature in Figure 3. Figure 5 shows theactivation energies as a function of the volume fraction of

Al2O3 in the LA mixtures. The reported activation energies ofLF mixtures32 are coplotted for comparison. The activationenergies of LA mixtures range from 0.44 to 0.56 eV below thephase transition temperature. Those values are much smallerthan the reported value of bulk LiBH4 (∼0.76 eV,orthorhombic phase).41 This result indicates that the interfacebetween LiBH4 and Al2O3 serves as a pathway for facilediffusion of Li ions. Such a decrease in activation energies wasconsistently found for LF mixtures. Nonetheless, the volumefraction of the insulating material at the activation energyminimum is smaller for the LA mixture than the LF mixture. Itis likely that the smaller particle size of Al2O3 than SiO2 shiftsthe volume fraction at the activation energy minimum to thelower value.At this point, we have demonstrated that the interface

between LiBH4 and Al2O3 significantly raises the ionicconductivity. However, identifying the physical and/orchemical origins which underlie such conductivity enhancementis challenging because it requires the structure analysis in thesubnanometer scale. NMR spectroscopy has been the mostcommon tool to probe the interface, but the focus so far hasbeen on the ionic mobility change instead of the chemicalbonding change.37,42−44 We employed NEXAFS spectroscopyto investigate the structural changes at the interface byanalyzing the chemical bonding of the constituting elements.Figure 6 shows B K-edge NEXAFS spectra of pure LiBH4 andLA mixtures with a different volume fraction of Al2O3. A verytiny peak appears around 194 eV in the TEY spectrum of pureLiBH4 while no distinct peak exists in the FY spectrum. As theamount of Al2O3 increases, the intensity of the peak at 194 eVincreases in both FY and TEY spectra, indicating that thebonding character of B has changed upon ball-milling withAl2O3. The peak at 194 eV is known to be the transition of B 1selectrons to the unoccupied π* state of [3]B with O (planarBO3),

45,46 and it applies to our study, as the intensity correlateswell with the amount of Al2O3. The tiny peak seen in the TEYspectrum of pure LiBH4 is likely to originate from the oxidizedsurface, which would be buried in the FY spectrum due to thereduced contribution by the surface. The new chemical bondsbetween B and O probed by the NEXAFS spectra reveal thatAl2O3 is not completely inert, which was not evident in theXRD pattern in Figure 1. Such an interfacial reaction possiblycontributes to the formation of a defective interface layer,thereby increasing the number of mobile Li ions and/orpromoting the migration of Li ions. The formation of B−O

Figure 4. (a) Comparison between the measured (solid symbols) and simulated (solid line) normalized conductivities for the LA mixtures at 25, 62,99, and 128 °C. (b) Arrhenius plots of the ionic conductivities of LA and LF interface layer calculated by a continuum percolation model.

Figure 5. Activation energy of Li conduction in ball-milled LA and LFmixtures as a function of volume fraction of oxides (Al2O3 or SiO2).

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bondings signifies the perturbation of chemical bondings at theinterface, but Li ions do not appear to migrate through theregion enriched with B−O bondings since the conductivity isnot simply proportional to the amount of B−O bondings whichincreases with the amount of Al2O3. Instead, the Li ionmigration is likely to occur through the interfacial LiBH4,possibly defective, directly contacting the region rich with B−Obondings. Therefore, even though the amount of B−Obondings increases, the conductivity inevitably starts todecrease when the amount of conducting phase, LiBH4,drops below a certain level.To probe any variations in the interfacial structures between

LA and LF mixtures, we compare the NEXAFS spectra of LA54 vol % and LF 55 vol % in Figure S4. When a similar volume% of oxide material is ball-milled with LiBH4, LA generates alarger amount of B−O bondings than LF does. As previouslydiscussed, the larger number of B−O bondings cannot bedirectly interpreted as higher conductivity. However, the resultat least indicates that γ-Al2O3 is more reactive or forms a largerinterface area compared to fumed silica, which indirectlysupports the higher conductivity of LA mixtures. Anotherinteresting point is the BO3 peak position (∼194 eV) shift tohigher energy in the case of LF. The shift to higher energy iscaused by shorter B−O bondings, different next nearest cations,etc.47 Therefore, the LA and LF mixtures have chemicallydistinct interfacial bondings which can result in the differentinterfacial conductivities.A recent first-principles study proposed that the low ionic

conductivity of orthorhombic LiBH4 is not due to lowermobility but due to lower carrier density.48 Thus, a highernumber of mobile Li ions, i.e. increased carrier concentration, atthe interface can significantly promote the ionic conductivity.Such a mechanism was manifested in the LiF−Li2CO3composite in which the carrier density is enhanced by themigration of Li interstitials from bulk LiF to the LiF/Li2CO3interface, leading to significantly improved ionic conductivity.49

In our case, the formation of the B−O bondings reflects apartial destruction of the inherent crystal structure of LiBH4 atthe interface, but we still need to clarify why mobile Li iondensity and/or mobility consequently becomes higher.Although the exact mechanism is subject to furtherinvestigation, our study highlights the interface engineering asan efficient route to increase ionic conductivity.

4. CONCLUSIONIn this study, we have characterized Li ion conductorscomposed of LiBH4 and Al2O3 and optimized theirconductivity by means of interface engineering. A mixture of

LiBH4−Al2O3 having different volume fractions of Al2O3 wassimply prepared by high energy ball-milling. TEM images andEELS element mapping data exhibit that LiBH4 and Al2O3 arewell mixed and create a large interface area during the ball-milling process. The ionic conductivity of mixture reaches anextremely high value of ∼2 × 10−4 S cm−1 at RT, which is 104

times higher than that of pure LiBH4. This enhancement inionic conductivity is probably due to the formation of a highlydefective interface between LiBH4 and Al2O3, supported by theformation of B−O bondings detected in the NEXAFS spectra,which would decrease the activation energy for Li migrationand/or increase the mobile Li ions. By analyzing theconductivity of the LA mixtures employing a continuumpercolation model, we have found that Al2O3 is even moreeffective in making a highly conducting interface than thepreviously investigated SiO2. The remarkable conductivityenhancement found in this study underscores the utility ofthe interface engineering as a simple and promising pathway forexploring new electrolytes for all-solid-state energy storage andconversion devices in the future.

■ ASSOCIATED CONTENT

*S Supporting InformationThe Supporting Information is available free of charge on theACS Publications website at DOI: 10.1021/acs.jpcc.7b08862.

Analytical details, Table S1, and Figures S1−S4 (PDF)

■ AUTHOR INFORMATION

Corresponding Author*E-mail: lee0su@kist.re.kr; Tel: +82-2-958-5412.

ORCID

Young-Su Lee: 0000-0002-3160-6633NotesThe authors declare no competing financial interest.

■ ACKNOWLEDGMENTS

This work was supported by the Innovation Fund Denmark viathe research project HyFill-Fast and by the TechnologyDevelopment Program to Solve Climate Changes of theNational Research Foundation (NRF) funded by the Ministryo f S c i e n c e a nd ICT , Ko r e a (G r an t Numbe r2015M1A2A2074688).

Figure 6. B K-edge NEXAFS spectra of pure LiBH4 and LA mixtures. Total electron yield and fluorescence yield for left and right panel, respectively.

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