2017 Volume 58 Issue 7 Pages 1063-1068
We have operated a 4V-class bulk-type, all-solid-state LiCoO2/Li battery at room temperature. The battery consisted of a Li4(BH4)3I complex hydride electrolyte as the electrolyte layer, and a 80Li2S 20P2S5 sulfide glass as an electrolyte in the positive electrode layer. The assembled battery exhibited a 92 mAh g−1 initial discharge capacity at 298 K and 0.1 C. The discharge capacity for the 20th cycle remained as high as 83 mAh g−1, corresponding to a capacity retention ratio of nearly 90%.
This Paper was Originally Published in Japanese in J. Japan Inst. Met. Mater. 80 (2016) 720–725.
Complex hydrides possess energy storage and conversion-related functions, including hydrogen storage,1,2) lithium storage,3,4) and fast ionic conductions.5–7) LiBH4, a typical complex hydride, experiences a phase transition from an orthorhombic structure (low-temperature phase, LT phase) to a hexagonal structure (high-temperature phase, HT phase) at elevated temperatures and around 390 K.8) At this point, a high lithium-ionic conductivity that exceeds log(σ/S cm−1) = −3 is realized in the HT phase of LiBH4,5) as shown in Fig. 1. The use of a complex hydride electrolyte for a battery assembly provides the following advantages: 1) use of materials that have a low gravimetric density (typically 1 g cm−3) is possible since light elements are chosen as constituents; 2) the transport number of lithium ions is close to unity since the materials are classified as electronic insulators; 3) complex hydrides allow for the establishment of a stable interface with a highly reactive metallic lithium electrode, which has the lowest voltage amongst the existing negative electrodes; 4) complex hydrides have high temperature durability since hydrogens in a crystal lattice are highly stabilized by a strong covalent bonding with center elements in a complex anion. The pyrolysis of complex hydrides is unlikely to occur in the temperature range considered for the battery operation; 6) the highly deformable nature of the complex hydrides allows for the preparation of an electrolyte compact merely by uniaxial pressing at room temperature, namely cold pressing. Introducing a robust interface between an electrode and an electrolyte is also easily performed by cold pressing.5–7)
In recent years, research and development efforts have been devoted to assembling bulk-type all-solid-state lithium rechargeable batteries that use complex hydride electrolytes, including the HT phase of LiBH4. LiBH4 has high reducing ability; this material is used as a reducing agent in chemical synthesis. For this reason, LiBH4 can form a stable interface with a highly reactive and low voltage lithium electrode.5–7) This high reducing ability of LiBH4 facilitates the reductive decomposition of positive electrodes, such as LiCoO2, which have oxidizing ability and high voltage, and thereby interferes with stable battery operation. To solve this issue, Takahashi and co-workers introduced an oxide-based solid electrolyte as a protection layer of a few tens of nanometers. This layer included Li3PO4 and avoided direct contact between the LiBH4 electrolyte and LiCoO2. As a result, repeated operation of the battery was successfully demonstrated.9,10) On the other hand, considering the high reducing ability of the complex hydride electrolyte, our research group has proposed a battery design principles that uses positive electrodes with a lower voltage and high capacity electrodes including TiS27,11) and elemental sulfur,12,13) combined with a lithium negative electrode. Repeated operation of the batteries based on this design concept has also been successfully demonstrated.
The operating temperature of a battery that uses the LiBH4 solid electrolyte is limited at temperatures higher than 390 K since the fast lithium-ionic conductivity of LiBH4 is realized in this temperature region.8) Hence, the developments of complex hydride-based solid electrolytes and battery design principles that allow for repeated operation of the all-solid-state battery across a wide-temperature range, including room temperature, is urgently required. Maekawa and co-workers revealed that the HT phase of LiBH4 is stabilized at reduced temperatures by the doping lithium halides such as LiCl, LiBr, and LiI. For instance, Li4(BH4)3I produces a high lithium-ionic conductivity of log(σ/S cm−1) = −4.7 at 300 K, as shown in Fig. 1.14) Sveinbjörnsson and co-workers assembled an all-solid-state battery that uses a Li4Ti5O12 porous thin-film electrode, a Li13/16(BH4)13/16I3/16 solid electrolyte, and a lithium negative electrode. They demonstrated repeated operation at 333 K.15) Yoshida and co-workers assembled a battery that uses a composite electrode containing a Li4(BH4)3I solid-electrolyte and a Li4Ti5O12 active material in an electrode layer. Battery operation across a wide-temperature range, i.e., 296 K to 423 K, was successfully demonstrated.16) An appropriate battery design principle that properly considers the electrochemical stability of an electrolyte would allow us to assemble all-solid-state batteries suitable for a room temperature operation.
As a typical sulfide-based solid electrolyte, a Li2S and P2S5 glass electrolyte is known to exhibit a high lithium-ionic conductivity (for instance, a 80Li2S 20P2S5 glass electrolyte having a room temperature lithium-ionic conductivity of log(σ/S cm−1) = −3.8, and an activation energy of 34 kJ mol−1)17). Owing to its highly deformable nature, this material allows for the assembly of a bulk-type all-solid-state battery merely by cold pressing. Repeated operation of a bulk-type all-solid-state battery that uses 4V-class LiCoO2 positive electrolyte at room temperature has been demonstrated.18) At the interface, which is comprised of a sulfide-based solid electrolyte and LiCoO2, high interface resistance is attributed to the formation of a space charge layer.19–22) In addition, with repeatd battery operation, a resistive layer is formed due to the mutual diffusion of constituent elements across the electrode and electrolyte interface.23) These issues are solved by coating solid electrolytes, including LiNbO3, at the LiCoO2 surface.
In this study, we assembled a 4V-class bulk-type all-solid-state lithium rechargeable battery that allows for room temperature operation. Considering the literature described above, both complex hydride and sulfide electrolytes are used for a battery assembly. A positive electrode active material, LiCoO2, whose surface was coated by LiNbO3 to a 20 nm thickness, and a lithium negative electrode are used. In the positive electrode layer, a mixture of the 80Li2S 20P2S5 solid electrolyte and the LiNbO3-coated LiCoO2 active material was used, while Li4(BH4)3I was employed as an electrolyte layer. Battery performance was evaluated at room temperature.
A Li4(BH4)3I electrolyte was synthesized via a mechanical ball milling technique and subsequent annealing.14) Powders of LiBH4 (≥90%, Sigma-Aldrich) and LiI (99.999%, Sigma-Aldrich) were used as starting materials. These powders were mixed using an agate pestle and mortal before the mixture was transferred to a 45 mL pot with 20 balls of 7 mm diameter. Mechanical ball milling was carried out at a rotation rate of 400 rpm for 5 h. From the resultant powder, no peaks other than the HT phase of LiBH4 appeared in the X-ray diffraction (XRD, X'pert PRO, PANalytical) patterns, and Raman spectra (Nicolet Almega XR, Thermo Fisher Scientific, Inc.).8)
To examine the oxidative stability of Li4(BH4)3I, a bulk-type all-solid-state lithium battery using a TiS2 positive electrode, the lithium negative electrode, and the Li4(BH4)3I electrolyte was assembled. Recently, the authors demonstrated repeated operation of over 300 cycles for a bulk-type all-solid-state lithium battery using the TiS2 positive electrode, the lithium negative electrode, and a LiBH4 electrolyte at 393 K.11) Hence, we consider that a similar battery configuration is appropriate for investigating the manner by which iodide-ions in the solid electrolyte affect the stability of the LiBH4 electrolyte and battery performance. In addition to this comparison, we also assembled a battery that uses a LiBH4 electrolyte instead of Li4(BH4)3I in the composite positive electrode, with which Li4(BH4)3I was used for the electrolyte layer.
TiS2, and Li4(BH4)3I or LiBH4, were weighted out to a 40:60 weight ratio and then mixed in an agate pestle and mortar. The resultant powders were used for the composite positive electrodes. The powders, 25 mg Li4(BH4)3I or 20 mg LiBH4, were transferred into an 8 mm diameter die and uniaxially pressed at 60 MPa. Subsequently, 6 mg composite positive electrodes were placed on the solid-electrolyte pellet still present in the die and further pressed at 240 MPa to produce a robust interface between the positive electrode and electrolyte layers. A lithium foil for use as a negative electrode was placed on the opposite side of the composite positive electrode layer. Battery tests were carried out at 393 K with a discharge–charge rate of 0.05 C (corresponding to a current density of 57 μA cm−2, 580 Battery Test System, Scribner Associates Inc.). The appearance of our battery and the configuration of the electrochemical system appeared have been given elsewhere.11–13)
For microstructure observation and element distribution analysis, a cross-section of the composite positive electrode was produced by a focused ion beam (FIB, FB2200, Hitachi High-Technologies Corp.) with Ga-ion beam radiation. For this purpose, a field-emission scanning electron microscopy (FE–SEM, SU9000, Hitachi High-Technologies Corp.) and an energy dispersive X-ray (EDX, Apollo XLT, Ametek, Inc.) were employed.
An 80Li2S 20P2S5 glass electrolyte was synthesized by the mechanical ball milling technique, as reported in a preceding paper.23) Powders of Li2S (98%, Sigma-Aldrich) and P2S5 (99%, Sigma-Aldrich) were mixed using an agate pestle and mortar. The resultant mixed powders were then transferred into a 45 mL zirconia pot with 160 zirconia balls of a 5 mm diameter. The mechanical milling was carried out at a rotation rate of 510 rpm for 10 h. Only a hollow pattern appeared in the XRD patterns of the resultant powders after mechanical ball milling; hence, we consider that glass electrolyte was obtained. The ionic conductivity of the prepared glass electrolyte compact was measured by a two-probe ac technique (3520-80 Chemical Impedance Meter, Hioki Corp.) in the temperature range of 298–423 K.
LiNbO3 was coated on to the surface of the LiCoO2 particles using a rolling fluidized coating machine (MP–01, Powrex Corp.)19,20) (this product is hereafter referred to as “LiNbO3-coated LiCoO2”). The powders of LiNbO3-coated LiCoO2, a Ketjen Black (KB) conductive additive, and the 80Li2S 20P2S5 glass electrolyte were weighed out to a 40:60:6 weight ratio, before being mixed and used for a composite positive electrode. Li4(BH4)3I electrolyte powder (101.1 mg) was transferred to a 10-mm-diameter die and then uniaxially pressed at 20 MPa and room temperature. Subsequently, 11.5 mg of composite positive electrode powder was placed on the electrolyte compact that was still present in the die, before being further pressed at 240 MPa. As a negative electrode, lithium foil was placed on the opposite side of the composite positive electrode. The battery test was carried out at 297 K and a charge–discharge rate of 0.1 C, corresponding to a current density of 73 μA cm−2, using an electroanalytical system, VMP3 (Bio-Logic Science Instruments). Microstructure observation by FE–SEM and element distribution analysis by EDX were carried out for the compact, which was comprised of the LiNbO3-coated LiCoO2, KB, and Li4(BH4)3I, to evaluate the thickness of the LiNbO3 layer at the LiCoO2 surface. A cross-section of this compact was produced by the FIB with Ga-ion beam radiation.
Despite the use of only cold pressing, a robust interface was introduced between the TiS2 active material and the Li4(BH4)3I electrolyte, as shown by the cross-sectional FE–SEM image, and by the iodine, titanium, and sulfur distributions of the composite positive electrode before the battery test (Fig. 2 (a)–(d)). Such an interface is expected to allow for a smooth charge transfer reaction during battery operation. This microstructure is similar to that of the composite positive electrode comprised of TiS2 and LiBH4.7,11) This suggests that high deformability is still maintained in the iodide-ion containing complex hydride, Li4(BH4)3I.
Microstructure and element distributions of the composite positive electrode layer, TiS2/Li4(BH4)3I. Before the battery test: (a) cross-sectional FE–SEM image, and distributions of (b) I, (c) Ti, and (d) S. After the battery test: (e) cross-sectional FE–SEM image, and distributions of (f) I, (g) Ti, and (h) S.
Battery reaction between TiS2 and lithium ions proceeds via the following electrochemical reaction,24,25)
| \[ {\rm TiS}_{2} + x{\rm Li}^{+} + x{\rm e}^{-} \rightleftarrows {\rm Li}_{x}{\rm TiS}_{2} \] | (1) |
Performance of the bulk-type all-solid-state batteries operated at 393 K and 0.05 C: (a) TiS2/Li4(BH4)3I | Li4(BH4)3I | Li, and (b) TiS2/LiBH4 | Li4(BH4)3I | Li.
To further examine how the iodide-ion in Li4(BH4)3I in the composite positive electrode affected battery performance at 393 K, we assembled a battery that used a LiBH4 solid electrolyte in the composite positive electrode, while Li4(BH4)3I was used for the electrolyte layer. For the battery configuration that used LiBH4 in the composite positive electrode, the initial discharge capacity was only 61 mAh g−1; this was due to the solid-state reaction between TiS2 and LiBH4 to form LixTiS2 (self-discharge). The second discharge capacity recovered to 218 mAh g−1, which is close to the theoretical capacity of TiS2 (239 mAh g−1). The initial charge capacity was 419 mAh g−1, which exceeded the theoretical capacity. This is due to the oxidative decomposition of unreacted LiBH4 of the self-discharge reaction before initial charge between TiS2 and LiBH4. This battery did not show any marked capacity reduction with repeated battery operation. Indeed, the 15th discharge capacity remained as high as 220 mAh g−1. This result suggests that the capacity fading of the bulk-type all-solid-state TiS2/Li4(BH4)3I | Li4(BH4)3I | Li battery, seen in Fig. 3 (a), is essentially attributable to interface instability between the TiS2 positive electrode and the Li4(BH4)3I solid electrolyte.
Microstructure observation and element distribution analysis of the composite positive electrode were carried out after the battery test to investigate how repeated battery operation affects the interface stability between TiS2 and Li4(BH4)3I. Figure 2 (e) shows a cross-sectional FE–SEM image of the positive electrode layer after the battery test, (f)–(h) show the distributions of iodine, titanium, and sulfur, respectively. After the battery test, iodine was concentrated at the interface between TiS2 and Li4(BH4)3I. The voltage drop became large with repeated battery operation, as the discharge–charge profiles shown in Fig. 3 (a). This suggests that the concentration of iodine was increased by the decomposition of Li4(BH4)3I due to the solid-state reaction with TiS2, resulting in the formation and precipitation of iodine-containing compounds and an increase in resistance. Hence, to realize repeated battery operation with low interface resistivity and high capacity, the use of a solid electrolyte that hinders such a side reaction (i.e., that leads to the precipitation and formation of iodine-containing species) would be beneficial.
Assembling an all-solid-state battery based on a sulfide-based solid electrolyte, and which has a lower internal resistance and allows for repeated battery operation, is possible when a 4V-class LiCoO2 electrode surface is coated by a solid electrolyte.19–23) Figure 4 (a) and (b) shows cross-sectional FE–SEM images of the composite comprised of LiNbO3-coated LiCoO2, KB, and Li4(BH4)3I. The distributions of cobalt, niobium, carbon, and iodine are shown in Fig. 4 (c)–(f). From these results, it is observed that LiCoO2 particles were homogeneously coated by LiNbO3 with a 20 nm thickness.
Microstructure of LiNbO3-coated LiCoO2 in the Li4(BH4)3I, KB, and LiNbO3-coated LiCoO3 compact: (a) and (b) cross-sectional FE–SEM image, and distributions of (c) Co, (d) Nb, (e) C, and (f) I.
The lithium-ionic conductivity of the 80Li2S 20P2S5 glass electrolyte synthesized in this study is shown in Fig. 1. The lithium-ionic conductivity at 298 K and activation energy were log(σ/S cm−1) = −4.0 and 29 kJ mol−1, respectively. These values relatively well with the values in the preceding work.17)
Electrochemical reaction of LiCoO2 proceeds as,26)
| \[ {\rm LiCoO}_{2} \rightleftarrows {\rm Li}_{1 - x} {\rm CoO}_{2} + x{\rm Li}^{+} + x{\rm e}^{-} \] | (2) |
Performance of the bulk-type all-solid-state battery, LiNbO3-coated LiCoO2/KB/80Li2S 20P2S5 | Li4(BH4)3I | Li, operated at 298 K and 0.05 C: (a) typical charge–discharge profiles, and (b) cycle performance.
In this study, we figured out that repeated battery operation is possible by a combined use of solid electrolytes, having different electrochemical stabilities, although the sole use of an electrolyte does not realize the repeated battery operation. Similar to this battery configuration, Kawaji and co-workers recently realized the operation of a high-temperature durable and sulfur-free 4V-class bulk-type all-solid-state battery by the combined use of a complex hydride electrolyte and an oxide electrolyte.27) A major remaining issue is achieving improved battery performance, including high capacity and low internal resistance, by further increasing the lithium-ionic conductivity of the complex hydride electrolyte at reduced temperatures, i.e., near to, or lower than, room temperature.
Recently, a solid electrolyte in which LiBH4 was dispersed in a Li2S–P2S5 glass matrix, realized a high lithium-ionic conductivity of log(σ/S cm−1) = −2.8 at 298 K.28) In addition, the crystalline phase of 90LiBH4–10P2S5, with a nominal composition, exhibited a lithium-ionic conductivity of log(σ/S cm−1) = −3.0 at 300 K.29) These materials do not include a phase transition, as seen in LiBH4, at least above room temperature; hence, the assembly of batteries that allowed for room temperature operation was possible. In other reports, high lithium30–33) and sodium30,32–35) ionic conductivities in complex hydrides containing cage-like cluster-anions, including [B12H12]2− and [CB11H12]−, have been discovered. Repeated battery operation of all-solid-state batteries by using some of the above materials has been successfully demonstrated.31,32) In complex hydride solid electrolytes, the cation transport rate is closely related to the reorientational dynamics of the complex anions, containing hydrogen.6,7,36–38) On the basis of this knowledge, diversification of the complex hydride electrolyte has rapidly progressed with regards to crystal structure and composition. To achieve better performance of bulk-type all-solid-state batteries, further exploration of this new-type of complex hydride solid electrolyte is highly desirable.
In this study, we assembled a 4V-class bulk-type all-solid-state lithium rechargeable battery by hybrid use of complex hydride and sulfide electrolytes. Repeated operation of this battery was successfully demonstrated at room temperature. For use as the lithium metal negative electrode, we chose Li4(BH4)3I electrolyte, which has a high reducing ability. A mixture of a LiCoO2 active material, whose surface was covered by LiNbO3 with a 20-nm thickness, a 80Li2S 20P2S5 glass solid electrolyte, and a KB conductive additive was used for the composite positive electrode. Our battery allowed for repeated operation at 298 K and a charge–discharge rate of 0.1 C. The initial and 20th discharge capacities of our battery were 92 mAh g−1 and 83 mAh g−1, respectively, and the capacity retention ratio was as high as 90%. This suggests that our battery design allows for stable battery operation at reduced temperature, e.g., 298 K.
The authors would like to thank Mr. K. Sato, Ms. H. Ohmiya, and Ms. N. Warifune for technical assistances. Financial supports from the Target Project 4 of WPI–AIMR, Tohoku University; Collaborative Research Center on Energy Materials, Tohoku University; JSPS KAKENHI Grant No. 25220911; and the Advanced Low Carbon Technology Research and Development Program (ALCA) from the Japan Science and Technology Agency (JST) are gratefully acknowledged.
