- Email: [email protected]
Electrochemical and interfacial behavior of all solid state batteries using Li10SnP2S12 solid electrolyte
Journal of Power Sources xxx (xxxx) xxx–xxx
Contents lists available at ScienceDirect
Journal of Power Sources journal homepage: www.elsevier.com/locate/jpowsour
Electrochemical and interfacial behavior of all solid state batteries using Li10SnP2S12 solid electrolyte Carolina Vinadoa, Shanyu Wanga, Yang Hec, Xingcheng Xiaob, Yun Lia, Chongmin Wangc, Jihui Yanga,∗ a
Materials Science and Engineering Department, University of Washington, Seattle, WA, 98195, USA Chemical Sciences and Materials Systems Lab, General Motors R&D Center, Warren, MI, 48090, USA c Pacific Northwest National Laboratory, Environmental Molecular Science Laboratory, Richland, WA, 99354, USA b
H I GH L IG H T S
utilizing Li SnP S solid electrolyte in all-solid-state batteries. • Successfully inter-diffusion between LCO and LSPS causes a poor cycling stability. • Atomic • Li NbO coating improves the performance of the LCO/LSPS solid state batteries. 10
3
2 12
4
A R T I C LE I N FO
A B S T R A C T
Keywords: Li10SnP2S12 All-solid-state battery Solid electrolyte Atomic layer deposition Interfacial stability
Thio-Lithium Superionic Conductor (Thio-LISICON) Li10GeP2S12 equivalent Li10SnP2S12 (LSPS) is comparable in ionic conductivity yet with a lower cost as an electrolyte for all solid-state batteries (ASSBs). ASSBs with LSPS solid electrolyte (SE), lithium-indium alloy anode, and LiCoO2 (LCO) cathode were successfully fabricated and their electrochemical performance at 60 °C was examined. Atomic layer deposition of Li3NbO4 on LCO was conducted to improve the interfacial stability. The Li3NbO4 coating effectively improves the cycle stability of the ASSB. Electrochemical impedance spectroscopy tests indicate a rapid growth of charge transfer resistance upon cycling for the cell with the uncoated LCO, primarily due to the surface instability and build-up of a space charge layer. However, the ASSBs with Li3NbO4 coated LCO show a more stable interface with a negligible impedance increase upon cycling, attributable to the buffering and passivating roles of the Li3NbO4 coating. The interfacial microstructure was analyzed to elucidate at the underlying reasons for the impedance increase and the pivotal role of the Li3NbO4 coating.
1. Introduction Lithium ion batteries (LIBs) are widely commercialized to power portable devices. Due to their high energy density, it is the most attractive option for hybrid electric vehicles, electric vehicles, and other portable applications [1]. Given the sensitive and widespread nature of these applications, high energy LIBs with superior reliability and safety are desirable. Most commercially available LIBs use flammable liquid electrolytes, which risks fire or explosions in the event of a failure [2]. From a safety point of view, replacing the flammable liquid electrolyte with a solid one would be desirable. An eligible solid electrolyte (SE) should have an ionic conductivity comparable to that of current liquid electrolytes, e.g., 10−3-10−2 S cm−1 [2]. Other benefits of SEs include larger electrochemical windows and excellent thermal stability, as well
∗
as the possibility of high power and energy density, such as when paired with a lithium metal anode [3]. The recent finding of Li10GeP2S12 (LGPS), with an outstanding ionic conductivity of 12 mS cm−1, comparable to liquid electrolytes, has made the use of sulfide-based SEs an appealing option for all solid-state batteries (ASSBs) [4,5]. The high cost of Ge, however, largely limits the practical use of LGPS. Creating isostructural analogues by substituting Ge with Si, Al, or Sn, has been suggested [6–8]. The Sn analogue Li10SnP2S12 (LSPS) is the most promising of the group [8]. The main hurdle for developing successful ASSBs, however, is in minimizing the interfacial impedances between the SE and the electrodes, especially for the cathode/SE interfaces. The interfacial impedances primarily dominate the rate capability and cycling stability of the ASSBs [1,4,9]. First, SEs must be chemically and electrochemically
Corresponding author. E-mail address: [email protected] (J. Yang).
https://doi.org/10.1016/j.jpowsour.2018.06.038 Received 6 March 2018; Received in revised form 8 June 2018; Accepted 9 June 2018 0378-7753/ © 2018 Elsevier B.V. All rights reserved.
Please cite this article as: Vinado, C., Journal of Power Sources (2018), https://doi.org/10.1016/j.jpowsour.2018.06.038
Journal of Power Sources xxx (xxxx) xxx–xxx
C. Vinado et al.
Li-In foil as the anode. The working electrode and LSPS were placed in a stainless steel die with a diameter of 10 mm and pressed together under 500 MPa for 20 min. The cathode/LSPS pellet was then dissembled from the die and placed into an ASSC. 60 mg of In foil (0.62 V vs Li+/Li [27], 99.99%, CMR-Direct, USA) and 10 mg of Li foil (99.9%, MTI Corporation, USA) were placed on the top of the SE and a pressure of 20 MPa was then applied. All cell preparation processes were carried in a glovebox (Lab Star, Mbraun, Germany, H2O < 0.5 ppm, O2 < 0.5 ppm). The ASSCs under a 20 MPa pressure were then galvanostatically charged/discharged using a Neware cycler in the voltage range of 2.0–3.7 V vs. Li-In at 60 °C. AC EIS measurements were performed in a potentiostat (Versastat4, Ametek Scientific Instruments, USA). EIS data of a cold-pressed SE pellet were collected at room temperature in a symmetrical cell using carbon foil as the electrodes in the frequency range of 1–106 Hz and with an amplitude of 5 mV. Similarly, EIS tests of the ASSCs in the frequency range of 0.01–106 Hz and an amplitude of 10 mV were performed at various cycling stages. Cyclic voltammetry (CV) was carried out on LSPS pellets sandwiched between a stainless steel (SS) plate and Li or In foil as electrodes at a scanning rate of 10 mV s−1 from 0 to 5 V. The crystal structure of the materials was investigated by powder Xray diffraction (XRD, Bruker D8 Discover with IμS 2D detector, Germany). The LSPS powders were first placed in a capillary tube and sealed with wax before exiting the glovebox. A field emission SEM (FESEM, FEI Sirion XL30, USA) equipped with an Oxford EDS was used to image the particle morphology and obtain elemental analysis and mapping. The microstructures and compositions of the cycled electrodes were analyzed by a scanning TEM equipped with a high-angle annular dark-field (HAADF) detector and an EDS at the Pacific Northwest National Laboratory (STEM/HAADF/EDS, FEI Titan 80–300 kV, USA). In order to reduce air- or moisture-exposure and avoid the unintended contaminations, the cycled cells were disassembled, the cathodes were scraped, and the cathode powders were placed on a TEM mesh in an Ar filled glovebox. The TEM holder was sealed and rapidly transferred into the TEM chamber with a minimal exposure.
stable against the electrode materials and able to form a stable and conductive interface (solid electrolyte interphase-like interface layer) [9–13]. Based on the theoretical calculations, sulfide-based SEs show narrow electrochemical windows, such as ∼1.8–2.3 V vs. Li+/Li for LGPS [14]. The narrow electrochemical windows of sulfides have been experimentally verified, although these had been largely overestimated by the conventional experimental method of using a Li/electrolyte/ inert metal semi-blocking electrode with limited contact areas [15]. In addition, due to the difference in Li+ chemical potential and resultant asymmetric Li+ migration, the depletion of Li+ on the sulfide side and the formation of a space charge layer largely increase the interfacial resistance [16]. To facilitate successful ASSBs, interfacial modifications on either electrodes or SE are necessary to mitigate the chemical diffusion and space charge layer, thus reducing the interfacial impedance and extending the electrochemical window beyond 4.2 V (vs. Li+/Li) [9–13,17–20]. To mitigate the interfacial resistance, the coating material should be a good ionic conductor, while thin and uniform. Typically, wet chemical methods such as sol-gel or spray coating, are used for creating surface coatings; these methods, however, leave non-uniform and thick coatings [21]. Alternatively, atomic layer deposition (ALD) boasts excellent conformity, atomic scale thickness control, and uniformity over large areas, emerging as a promising process for battery electrode coating [9,22]. The sequential and self-limiting nature of ALD allows for ultra-thin and pinhole free coatings to be applied. These subnanometer coating layers have the added benefits of lessened mechanical stresses due to the crystal mismatch and reduced thickness constraints as far as the electron and lithium ion conductivities of the coating material are concerned [21]. To evaluate the potential of LSPS as an SE and study its interfacial behavior with typical cathode materials, ASSBs with LSPS SE, Li/In anode, pristine and ALD Li3NbO4-coated LiCoO2 cathodes were fabricated. Li3NbO4 was chosen because of its reasonable ionic conductivity and excellent chemical stability against both layered oxide cathodes and LSPS, thus effectively protecting the cathode from undesirable side reactions with the electrolytes [23–26]. The electrochemical performance and interfacial characteristics of the ASSBs using pristine and coated LiCoO2 were systematically explored by galvanostatic cycling, AC electrochemical impedance spectroscopy (EIS), scanning electron microscope (SEM), transmission electron microscope (TEM), energy dispersive X-ray spectroscopy (EDS), etc. The study shows that the ALD coating largely improves the electrochemical performance of ASSBs with the LSPS SE, in large part due to the improved interfacial compatibility and decreased charge transfer impedance.
3. Results and discussion 3.1. Characterization of the solid electrolyte and cathode materials Fig. 1a shows the crystal structure of LSPS, taken from the online database Materials Project [28]. LSPS crystalizes in a tetragonal structure with a space group P42/mc. The lattice parameters are a = 8.854 Å and c = 12.851 Å. The three-dimensional percolating structure ensures high Li+ conductivity. To verify the crystal structure of the as-received LSPS, powder XRD was carried out, as shown in Fig. 1b. Based on the above crystal structure, the standard XRD pattern of LSPS was also simulated using the Materials Studio's software, also shown in Fig. 1b. Our simulated XRD matches well with the result presented in Materials Project and those given by Bron, Kuhn, and Tarhouchi et al. [7,29,30]. As shown, the experimental XRD can be well indexed to the simulated XRD, indicating good phase purity of the LSPS powders. In addition, the sharp and well-defined diffraction peaks indicate good crystallinity of the SE. Fig. 1c shows the Nyquist plot for an LSPS pellet cold pressed at 500 MPa. Similar to the previous work done by Bron et al. [7], we observed grain resistance and relatively small grain boundary contribution, even without heat treatment. The Nyquist plot can be well fitted by the equivalent circuit shown in the inset of Fig. 1b. Rb, corresponding to the high-frequency intercept with the real axis, is the bulk resistance; the inductance segment L is mainly due to the cabling and leads of the set up [31]; the Rgb|CPEgb combination, showing as the semicircle in the mid-low frequency range, corresponds to the ionic resistance and capacitance of the grain boundaries in the LSPS pellet [8]. Based on the bulk resistance and geometry of the pellet, a total
2. Experimental ALD of ultrathin lithium niobium oxide coatings were deposited using a Savannah 100 Atomic Layer Deposition system (Cambridge Nanotech Inc., USA), with lithium tert-butoxide (LiOtBu, (CH3)3COLi, Sigma Aldrich, USA) and niobium ethoxide [C10H25NbO5, Sigma Aldrich] as the precursors. The distilled water was used as the oxidant and Argon (99.99%) as the carrier gas. The pulse, exposure, and purge times for all precursors are 0.1 s, 10 s, and 20 s, respectively, in each sub-cycle. The heating temperature was 160 °C for both Li and Nb precursors, and 80 °C for water. The deposition temperature was around 200 °C for the LiCoO2 powder samples (LCO, 99.8%, Sigma Aldrich, USA). The electrochemical properties LCO/LSPS/Li-In ASSBs were investigated using a homemade all-solid-state cells shown in Fig. S1 (Supporting Information, SI). Electrode mixtures consisted of pristine LCO or Li3NbO4 coated LCO (c-LCO) powders, and LSPS powders (Nanomyte SSE-10, NEI Corporation, USA) in a weight ratio of 70:30. They were prepared by hand grinding in a mortar for 30 min. Twoelectrode all solid state cells (ASSCs) were fabricated using ∼ 12.7 mg of electrode mixture as the cathode, 80 mg LSPS as the separator, and 2
Journal of Power Sources xxx (xxxx) xxx–xxx
C. Vinado et al.
Fig. 1. (A) Crystal structure of LSPS. [SnS4] (Orange) and [PS4] (Aqua) tetrahedra are shown. (b) Experimental and simulated powder X-ray diffraction (XRD) of LSPS. (c) Nyquist plot of the experimental (Black Square) and fit (Red line) of a LSPS pellet in a C/LSPS/C cell at room temperature. (d) Cyclic voltammograms of the Li/LSPS/SS and In/ LSPS/SS for the electrochemical stability window evaluation of the SE vs. Li & In between −0.5 and 5 V vs. Li+/Li at room temperature. Scanning electron microscopy images of the (e) LSPS, (f) uncoated LCO, and (g) c-LCO powders. (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.)
ionic conductivity of ∼3 mS cm−1 at room temperature can be obtained for our cold-pressed LSPS sample, comparable to the literature values [7,8,30,32,33]. To select an appropriate anode material, cyclic voltammetry was done on a cold pressed LSPS pellet. Fig. 1d shows the cyclic voltammograms of LSPS to evaluate the electrochemical stability window in Li/LSPS/SS and In/LSPS/SS asymmetric cells. The red line is the CV curve for the Li/LSPS/SS setup. A large cathodic peak, due to the deposition of metallic lithium (Li+ + e−→ Li), and a slight anodic peak associated with the stripping of metallic lithium (Li → Li+ + e−), are observed at around 0 V, similar to those of related systems [30,32,33]. On subsequent cycles, the peak's currents gradually decrease (see details: Fig. S2). It is well known that sulfide-based SEs are unstable against Li metal, which can cause a reduction of the SE upon contact with Li metal [30]. This is verified by the significant loss of current signal on the following cycles, further ascribed to the largely increased resistance from the interfacial reaction products. For the In/LSPS/SS setup, however, no noticeable change in the cathodic and anodic peaks can be observed during the voltage sweep (black line in Fig. 1d), implying no redox occurring for the In/LSPS interface. Thus, Li-In alloys with a high oxidation potential of ∼0.62 V vs. Li+/Li and good ductility were chosen for our experiments to ensure good contact with LSPS and bypass any confounding factors arising from parasitic reactions during our electrochemical tests. Fig. 1e shows a SEM image of the LSPS SE as received. The particles are irregularly shaped oblongs with sizes of 1–10 μm and significant aggregations. EDS (Fig. S3) shows a fairly uniform elemental distribution, consistent with the single phase XRD pattern. Similarly, Fig. 1f and g shows the SEM images of the pristine and Li3NbO4-coated LCO, respectively, both of which have a chunk or irregular-shape morphology
and a wide size distribution of 1–20 μm. EDS mapping of the coated sample shows a fairly even elemental distribution (Fig S4 in SI), consistent with its single phase XRD (Fig.S5 in SI). The coating layer is undetectable under SEM/EDS or XRD, primarily due to the thin nature and low volume percentage content of the coating (only 5 ALD cycles and ∼0.5 nm thick). 3.2. Electrochemical performance of the all-solid-state cells In order to study the electrochemical performance of the SE and its compatibility with LCO, ASSBs with pristine and c-LCO were fabricated and tested. Fig. 2a shows a comparison of the cycling stability between LCO and c-LCO. First, LCO shows a low discharge capacity of ∼105 mA h/g at C/10 and 60 °C, much lower than the value achieved in liquid cells (∼145 mA h/g at an upper cut-off of 4.3 V vs. Li+/Li) [32]. More importantly, the initial Coulombic efficiency is only 82%, also much lower than that of liquid cells (98%) [34]; similarly, the system in general possesses other limitations when compared to liquid electrolyte systems, such as added interfacial resistances, impaired kinetics under certain circumstances, and a thicker electrolyte layer. The significant initial capacity loss indicates an interfacial instability between LSPS and LCO that consumes Li in order to form a solid electrolyte interphase (SEI). The large voltage polarization of the LCO cell also indicates its large interfacial impedance. The LCO cell activates its discharge capacity to ∼110 mA h/g in the 2nd cycle and then experiences a severe capacity decay, only retaining ∼60% initial capacity after 70 cycles. This large capacity decay is accompanied by a large voltage decline. For instance, the average discharge voltage increases from 3.24 V for the 1st cycle to 3.28 V on the 10th cycle and gradually decreases to 3.15 V by the 50th cycle., as seen in Fig. 2b. The large 3
Journal of Power Sources xxx (xxxx) xxx–xxx
C. Vinado et al.
Fig. 2. (A) Cycling performance of LCO/LSPS/Li-In and c-LCO/LSPS/Li-In between 2.0 and 3.7 V vs. Li-In (2.62–4.32 V vs Li+/Li) at C/10 (0.13 mA/cm2) and 60 °C. Voltage profiles of (b) LCO/LSPS/Li-In and (c) c-LCO/LSPS/Li-In.
interdiffusion are allowed to occur for a longer time when using a low C-rate, resulting in a severally deteriorated performance. The cell shows even lower capacities at higher C-rates: ∼20 mA h/g at C/10, 10 mA h/ g at C/5, and no observable capacity at C/3 or higher. In addition, the LCO cell only retains 15 mA h/g discharge capacity when cycled at C/ 10 again, demonstrating a failure of the cell. Conversely, the c-LCO cell has a much better rate capability, showing a discharge capacity of ∼120 mA h/g at C/20, 115 mA h/g at C/10, 95 mA h/g at C/5, 70 mA h/g at C/3, and 20 mA h/g at 1C, as shown in Fig. 3a. After the high-rate cycles, the c-LCO cell can recover ∼95% initial capacity at C/ 10, corroborating its much better cyclability. A comparison of voltage profiles, shown in Fig. 3b, clearly indicate higher discharge voltages and lower overpotentials of the c-LCO cell, as compared with the LCO cell. The noticeably improved cycling stability and rate capability of the c-LCO cell are primarily originated from the reduced interfacial resistance and improved interfacial stability, which will be discussed in more details below.
voltage decline is a consequence of a radically increasing LCO/LSPS interfacial resistance, contributing to the severe capacity decay, together with the active material loss, as LCO is consumed by the reactions with LSPS, which will be discussed later. The c-LCO cell not only shows a higher discharge specific capacity but also better cycling stability. The initial specific discharge capacity and Coulombic efficiency of the c-LCO cell are ∼118 mA h/g and ∼87%, respectively. Its capacity gradually activates to ∼127 mA h/g by the 5th cycle, consistent with the decreasing overpotential and interfacial impedance of the cell. The activation process is presumably originated from the interface activation (improved particle contact, SEI formation, void filling, etc.). Since Li3NbO4 is stable against both LCO and LSPS, the coating largely avoids the direct physical contact of LCO and LSPS and thus mitigates their chemical reactions, protecting the interface in the following cycles. In addition, due to the buffering effect of the coating, the space charge layer and resultant resistance are largely reduced [16]. These two effects contribute to stable LCO/Li3NbO4/ LSPS interfaces and thus a much better cycling stability. The c-LCO cell can retain ∼85% initial capacity after 70 C/10 cycles, outperforming the LCO cell. More importantly, the average discharge voltage of the cLCO cell initially increases from 3.14 V for the 1st cycle to 3.42 V for the 10th cycle, and slowly declines to 3.24 V for the 50th cycle, as shown on Fig. 2c. The initial capacities of the uncoated and coated LCO, as well as their cycling behavior is well consistent with those of other LCO ASSBs [12]. Fig. 3a shows the rate capability of the LCO and c-LCO ASSBs. Due to severe interfacial instability and thus high resistance, pristine LCO shows poor rate capability and cycling stability, as shown in Fig. 3a. The LCO cell only shows a low initial capacity of ∼77 mA h/g and loses almost 75% of its capacity in first 5 cycles at C/20. The extremely low initial capacity may be related to severe side reactions of LCO/LSPS at the slow charge process (C/20), as the chemical reactions and elemental
3.3. Electrochemical impedance spectroscopy study As demonstrated above, surface coating to physically separate LCO and LSPS is indispensable for employing this highly conductive and technically feasible SE. The improved interfacial stability is evidenced by the enhanced cycling stability and rate capability, which can be further clarified by monitoring the impedance evolution with cycling. Fig. 4a and b depict the room temperature Nyquist plots of the LCO and c-LCO cells at different cycling stages. The Nyquist plots could be well fitted by the equivalent circuit shown in the inset of Fig. 4a, which contains the serial combination of internal resistance RS (high frequency intercept with real axis), interfacial resistance Rsei (high frequency semicircle, ∼100 kHz), charge transfer resistance Rct (low frequency semicircle, ∼500–600 Hz), and Warburg element representing 4
Journal of Power Sources xxx (xxxx) xxx–xxx
C. Vinado et al.
Fig. 3. (A) Rate capability of LCO/LSPS/Li-In and c-LCO/LSPS/Li-In between 2.0 and 3.7 V vs. Li-In (2.62–4.32 V vs. Li+/Li) at 60 °C. (b) Voltage profiles at different rates of LCO/LSPS/Li-In (dashed lines) and c-LCO/LSPS/Li-In (solid lines).
indicates the interfacial instability between LCO and LSPS, due to chemical reactions, mutual elemental diffusions, etc. The interfacial analyses will be presented below to clarify this. The thin Li3NbO4 coating does improve the interfacial stability of LCO/LSPS, although it cannot completely get rid of the Rsei. Similar effects can be found for ASSBs with other coated materials [9–13,18,20]. Further modification of the coating thickness is required to mitigate the SEI formation, and it is in our plans for future work. Large charge transfer resistances Rct, ∼13700 Ω cm and 7000 Ω cm for the LCO and c-LCO cells, respectively, can be observed which can be attributed to the resistive surface layers originated from surface contamination of LSPS and the reactivity of LCO/LSPS. The resistive surface layers not only impede Li+ transfer but also limit the electron transfer across active particle surfaces. Similar to the Rsei, Rct also experiences an activation process during the first several cycles; upon assembly of the cell, a chemically inactive surface layer forms; once the cell is cycled, this layer is destroyed and modified such that it is less resistive and facilitates the charge transfer process [26]. This is referred as the ‘activation process’, and the reason for the initial decrease in impedance. After the fifth cycle, Rct reduces noticeably to 1700 and 900 Ω cm for the LCO and c-LCO cells, respectively. The decrease in Rct should share the same origin as that of Rsei, associated with the removal/formation of the passivation layer on the particle surfaces and/ or increased contact area. In addition, the lower Rct of the c-LCO cell is mainly ascribed to the surface protection by the coating, which largely
the solid-state ion diffusion (low frequency tail), as well as the respective constant phase elements (CPE) associated with Rsei and Rct [17–20]. The fitting results of RS, Rsei, and Rct are summarized in Fig. 4c. The internal resistance of the as-prepared cells, including the resistances from cathode, anode, separator, current collectors, and leads, are ∼560 Ω cm for the LCO cell and ∼790 Ω cm for the c-LCO cell; the higher RS of the c-LCO cell is mainly due to the electronically insulating coating. The significant RS drop after the 1st cycle for both cells is primarily attributed to improved contact and the activation of particles after high-temperature cycling. The c-LCO cell keeps its low RS (210 Ω cm) after 10 cycles, while the LCO cell shows largely increased RS (680 Ω cm), consistent with its interfacial instability. Due to the air- and moisture-sensitivity of LSPS, the LSPS particle surface is inevitably contaminated and forms a passivating layer. Although LCO is stable in air, its surface could also absorb trace amounts of H2O, which further contaminates the LSPS surface when mixing and assembling them together. These facts could account for the large Rsei of the as-assembled LCO (∼2500 Ω cm) and c-LCO cells (∼3530 Ω cm). Interestingly, after one charge at 60 °C Rsei of both cells almost disappears (∼200 Ω cm). The exact origin of this phenomenon, however, needs further experimental and theoretical efforts. Rsei increases slightly to ∼450 Ω cm after 5 cycles, however, it jumps to large values of ∼6740 Ω cm and ∼2590 Ω cm for the LCO and c-LCO cells after 10 cycles, respectively. The significant increase in Rsei clearly
Fig. 4. Room Temperature Nyquist plots of (a) LCO/LSPS/Li-In and (b) c-LCO/LSPS/Li-In at different cycling stages (as-prepared and after 1st, 5th, and 10th charge). (c) Summary of Rs, Rsei, and Rct with cycling derived from EIS fitting. 5
Journal of Power Sources xxx (xxxx) xxx–xxx
C. Vinado et al.
electron transport [37,38], resulting in the high resistance and considerable overpotential of the LCO/LSPS/Li-In cell. Fig. 5b shows a TEM image of the c-LCO/LSPS interface after 10 C/ 10 cycles at 60 °C. Due to the ultrathin coating and beam-sensitivity of Li-ion conductors, Li3NbO4 is undetectable. A clear and narrow interface, however, can be observed, as compared with that of the LCO/LSPS interface. The EDX line scan of the Co, Sn, and P elements across the interface shows a thin inter-diffusion length of ∼4 nm. Co diffusion into LSPS was mostly blocked, and P and Sn diffuse ∼4 nm into c-LCO. It is obvious that the ultrathin coating layer largely impedes the atomic inter-diffusion, resulting in much thinner “damage layer” and hence lower interfacial resistance compared with uncoated particles. The Li3NbO4 coating, as suggested [39], also presumably alleviates the space charge layer formation, further reducing the Li+ transfer resistance. As shown in Fig. 4c, the c-LCO cell still suffers large impedance growth after 10 cycles, probably due to the ultrathin coating in this work. Further optimization of the coating thickness is required to optimize the interface behavior and cycling performance. Based on the above results, in order to use the highly conductive and cost-effective LSPS in ASSBs, surface protection of the cathode materials is necessary, at least for the layered cathodes, i.e., LCO, NMC (LiNi1-x-yMnxCoyO2), NCA (LiNi0.8Co0.15Al0.05O2). The high reactivity of sulfide-based SEs mainly originates from the high malleability of S2−, which also gives rise to its merits: high ionic conductivity and good deformability, as compared with oxide-based SEs. Here we have successfully demonstrated that an ALD coating with ultrathin oxide-based Li-ion conductors is effective at not only protecting the active-material/ SE interface, but also reducing interfacial resistance, thus mitigating the cell impedance and its build-up with cycling.
prevents the formation of a thick SEI [3,9,32]. The activation process, including the reduction of all three resistances, agrees well with the reduction of overpotential and increased specific capacity in the initial several cycles, as shown in Fig. 2. The capacity and impedance activation process is widely observed in ASSBs [35] and liquid cells with surface sensitive active materials, such as NMC811 (LiNi0.8Mn0.1Co0.1O2) [36]. After 10 cycles, Rct increases significantly to ∼9800 Ω cm and 4200 Ω cm for the LCO and c-LCO cells, respectively. It is obvious that the SEI and charge transfer resistances dominate the cell performances, including the cycling stability and rate capability. After 10 cycles, the LCO cell shows a large total resistance of ∼17000 Ω cm, much larger than that of the c-LCO cell (∼7000 Ω cm). The resistance of the LCO cell will keep increasing without forming stable interfaces, and finally terminate the cell due to a strong voltage polarization. Surface coating with Li3NbO4 largely improves the interfacial stability, and thus the cycling stability and rate capability, as demonstrated in other studies [3,9,13,17,34].
3.4. TEM analysis of LCO/LSPS and c-LCO/LSPS interfaces To gain better understanding of the LCO/LSPS interfacial behavior and the effects of the surface coating, TEM observations of the composite cathodes were carried out to directly visualize the LCO/LSPS interface. The ASSBs were disassembled after 10 C/10 cycles at 60 °C in an argon-filled glove box. Figure 5a shows a TEM image of an LCO/ LSPS interface, and a thick passivating layer with gradually varying contrast can be observed in between. An EDX line scan across the interface confirms the interface thickness to be of ∼18 nm. The variations of Co, Sn, and P elements at the interface show that the mutual diffusion length is ∼5 nm on both sides. Particularly, Sn is largely depleted in the LSPS side, indicating that Sn deeply migrates into LCO. A high-resolution TEM image, shown in Fig. S6 (SI), indicates that the interface seems to be composed of crystalline nanoparticles, which are most likely cobalt and/or tin sulfides. Similar results have been found for Li2S-P2S5 SEs interphases [13,18]. The sulfides, which presumably make up this interfacial layer, are highly resistive for both Li+ and
4. Conclusions All solid-state batteries with LiCoO2 cathodes, Li10SnP2S12 solid electrolyte, and Li-In anode were fabricated to evaluate the electrochemical performance of low-cost Li10SnP2S12 and its interfacial behavior with LiCoO2. Poor cycling stability and rate capability can be
Fig. 5. TEM images and corresponding EDX line scans of (a) LCO/LSPS and (b) c-LCO/LSPS interfaces after 10 C/10 cycles at 60 °C. 6
Journal of Power Sources xxx (xxxx) xxx–xxx
C. Vinado et al.
observed for the cell with pristine LiCoO2, primarily due to the instability of the LCO/LSPS interface. Considerable atomic inter-diffusion causes a thick interphase layer and markedly increases the interfacial resistance, resulting in large voltage polarizations and poor cyclability. Surface coating of ultrathin Li3NbO4 on LiCoO2 by atomic layer deposition largely improves the cycling and rate performance of the ASSBs. The enhanced electrochemical performance is further traced back to the decreased charge transfer and interfacial resistances, originated from the mitigation of atomic inter-diffusion and the formation of resistive interphases. Our study also demonstrated that ALD coating with thin layer oxide-based Li-ion conductors is imperative and effective for the utilization of sulfide-based SEs for ASSBs.
[14]
[15] [16]
[17]
[18]
[19]
Acknowledgements [20]
CMW thanks the support of the Assistant Secretary for Energy Efficiency and Renewable Energy, Office of Vehicle Technologies of the U.S. Department of Energy under Contract No. DE-AC02-05CH11231, Subcontract No. 18769 and No. 6951379 under the Advanced Battery Materials Research (BMR) program. The microscopic analysis in this work was conducted in the William R. Wiley Environmental Molecular Sciences Laboratory (EMSL), a national scientific user facility sponsored by DOE's Office of Biological and Environmental Research and located at PNNL. PNNL is operated by Battelle for the Department of Energy under Contract DE-AC05-76RLO1830. X.X also acknowledges the support by the Assistant Secretary for Energy Efficiency and Renewable Energy, Vehicle Technologies Office of the U.S. Department of Energy under Contract No. Award Number DE-EE0007787 under the Battery Material Research (BMR) Program.
[21]
[22]
[23]
[24]
[25]
[26]
Appendix A. Supplementary data Supplementary data related to this article can be found at http://dx. doi.org/10.1016/j.jpowsour.2018.06.038.
[27]
References
[28]
[1] B. Wu, S. Wang, W.J. Evans IV, D.Z. Deng, J. Yang, J. Xiao, Interfacial behaviours between lithium ion conductors and electrode materials in various battery systems, J. Mater. Chem. 4 (2016) 15266–15280, http://dx.doi.org/10.1039/C6TA05439K. [2] J.B. Goodenough, Y. Kim, Challenges for rechargeable Li batteries, Chem. Mater. 22 (2010) 587–603, http://dx.doi.org/10.1021/cm901452z. [3] A. Sakuda, A. Hayashi, T. Ohtomo, S. Hama, M. Tatsumisago, All-solid-state lithium secondary batteries using LiCoO2 particles with pulsed laser deposition coatings of Li2S–P2S5 solid electrolytes, J. Power Sources 196 (2011) 6735–6741, http://dx.doi.org/ 10.1016/j.jpowsour.2010.10.103. [4] N. Kamaya, K. Homma, Y. Yamakawa, M. Hirayama, R. Kanno, M. Yonemura, T. Kamiyama, Y. Kato, S. Hama, K. Kawamoto, A. Mitsui, A lithium superionic conductor, Nat. Mater. 10 (2011) 682–686, http://dx.doi.org/10.1038/nmat3066. [5] Y. Kato, S. Hori, T. Saito, K. Suzuki, M. Hirayama, A. Mitsui, M. Yonemura, H. Iba, R. Kanno, High-power all-solid-state batteries using sulfide superionic conductors, Nat. Energy (2016) 16030, http://dx.doi.org/10.1038/nenergy.2016.30. [6] S.P. Ong, Y. Mo, W.D. Richards, L. Miara, H.S. Lee, G. Ceder, Phase stability, electrochemical stability and ionic conductivity of the Li10 ± 1MP2X12 (M = Ge, Si, Sn, Al or P, and X = O, S or Se) family of superionic conductors, Energy Environ. Sci. 6 (2013) 148–156, http://dx.doi.org/10.1039/C2EE23355J. [7] P. Bron, S. Johansson, K. Zick, J. Schmedt auf der Günne, S. Dehnen, B. Roling, Li10SnP2S12: an affordable lithium superionic conductor, J. Am. Chem. Soc. 135 (2013) 15694–15697, http://dx.doi.org/10.1021/ja407393y. [8] P. Bron, S. Dehnen, B. Roling, Li10Si0.3Sn0.7P2S12 – a low-cost and low-grain-boundaryresistance lithium superionic conductor, J. Power Sources 329 (2016) 530–535, http:// dx.doi.org/10.1016/j.jpowsour.2016.08.115. [9] J.H. Woo, J.E. Trevey, A.S. Cavanagh, Y.S. Choi, S.C. Kim, S.M. George, K.H. Oh, S. Lee, Nanoscale Interface Modification of LiCoO2 by Al2O3 Atomic Layer Deposition for Solidstate Li Batteries, vol. 159, (2012), pp. 1120–1124, http://dx.doi.org/10.1149/2. 085207jes. [10] N. Ohta, K. Takada, L. Zhang, R. Ma, M. Osada, T. Sasaki, Enhancement of the high-rate capability of solid-state lithium batteries by nanoscale interfacial modification, Adv. Mater. 18 (2006) 2226–2229, http://dx.doi.org/10.1002/adma.200502604. [11] K. Takada, T. Ohno, N. Ohta, T. Ohnishi, Y. Tanaka, Positive and negative aspects of interfaces in solid-state batteries, ACS Energy Lett. 3 (2018) 98–103, http://dx.doi.org/ 10.1021/acsenergylett.7b01105. [12] A. Sakuda, H. Kitaura, A. Hayashi, K. Tadanaga, M. Tatsumisago, J. Power Sources 189 (2009) 527–530. [13] A. Sakuda, H. Kitaura, A. Hayashi, K. Tadanaga, M. Tatsumisago, All-solid-state lithium
[29]
[30]
[31]
[32]
[33]
[34]
[35]
[36]
[37]
[38]
[39]
7
secondary batteries with oxide-coated LiCoO2 electrode and Li2S–P2S5 electrolyte, J. Power Sources 189 (2009) 527–530, http://dx.doi.org/10.1016/j.jpowsour.2008.10.129. Y. Zhu, X. He, Y. Mo, First principles study on electrochemical and chemical stability of solid electrolyte–electrode interfaces in all-solid-state Li-ion batteries, J. Mater. Chem. 4 (2016) 3253–3266, http://dx.doi.org/10.1039/C5TA08574H. F. Han, T. Gao, Y. Zhu, K.J. Gaskell, C. Wang, A battery made from a single material, Adv. Mater. 27 (2015) 3473–3483, http://dx.doi.org/10.1002/adma.201500180. J. Haruyama, K. Sodeyama, L. Han, K. Takada, Y. Tateyama, Space–charge layer effect at interface between oxide cathode and sulfide electrolyte in all-solid-state lithium-ion battery, Chem. Mater. 26 (2014) 4248–4255, http://dx.doi.org/10.1021/cm5016959. Y. Iriyama, H. Kurita, I. Yamada, T. Abe, Z. Ogumi, Effects of surface modification by MgO on interfacial reactions of lithium cobalt oxide thin film electrode, J. Power Sources 137 (2004) 111–116, http://dx.doi.org/10.1016/j.jpowsour.2004.05.029. Y. Huang, J. Chen, J. Ni, H. Zhou, X. Zhang, A modified ZrO2-coating process to improve electrochemical performance of Li(Ni1/3Co1/3Mn1/3)O2, J. Power Sources 188 (2009) 538–545, http://dx.doi.org/10.1016/j.jpowsour.2008.12.037. Y. Huang, J. Chen, F. Cheng, W. Wan, W. Liu, H. Zhou, X. Zhang, A modified Al2O3 coating process to enhance the electrochemical performance of Li(Ni1/3Co1/3Mn1/3)O2 and its comparison with traditional Al2O3 coating process, J. Power Sources 195 (2010) 8267–8274, http://dx.doi.org/10.1016/j.jpowsour.2010.07.021. J.-H. Shim, K.-S. Lee, A. Missyul, J. Lee, B. Linn, E.C. Lee, S. Lee, Characterization of spinel LixCo2O4 -coated LiCoO2 prepared with post-thermal treatment as a cathode material for lithium ion batteries, Chem. Mater. 27 (2015) 3273–3279, http://dx.doi.org/10. 1021/acs.chemmater.5b00159. Z. Chen, Y. Qin, K. Amine, Y.-K. Sun, Role of surface coating on cathode materials for lithium-ion batteries, J. Mater. Chem. 20 (2010) 7606, http://dx.doi.org/10.1039/ c0jm00154f. H. Kaneda, Y. Koshika, T. Nakamura, H. Nagata, R. Ushio, K. Mori, Improving the cycling performance and thermal stability of LiNi0.6Co0.2Mn0.2O2 cathode materials by Nb-doping and surface modification, Int. J. Electrochem. Sci. 12 (2017) 4640–4653, http://dx.doi. org/10.20964/2017.06.19. K. Takada, N. Ohta, L. Zhang, K. Fukuda, I. Sakaguchi, R. Ma, M. Osada, T. Sasaki, Interfacial modification for high-power solid-state lithium batteries, Solid State Ionics 179 (2008) 1333–1337, http://dx.doi.org/10.1016/j.ssi.2008.02.017. B. Ahmed, C. Xia, H.N. Alshareef, Electrode surface engineering by atomic layer deposition: a promising pathway toward better energy storage, Nano Today 11 (2016) 250–271, http://dx.doi.org/10.1016/j.nantod.2016.04.004. K. Takada, N. Aotani, K. Iwamoto, S. Kondo, Solid state lithium battery with oxysulfide glass, Solid State Ionics 86 (1996) 877–882, http://dx.doi.org/10.1088/1751-8113/44/ 8/085201. S. Liu, X. Chen, J. Zhao, J. Su, C. Zhang, T. Huang, J. Wu, A. Yu, Uncovering the role of Nb modification in improving the structure stability and electrochemical performance of LiNi0.6Co0.2Mn0.2O2 cathode charged at higher voltage of 4.5 V, J. Power Sources 374 (2018) 149–157, http://dx.doi.org/10.1016/j.jpowsour.2017.11.037. N. Ohta, K. Takada, I. Sakaguchi, L. Zhang, R. Ma, K. Fukuda, M. Osada, T. Sasaki, LiNbO3-coated LiCoO2 as cathode material for all solid-state lithium secondary batteries, Electrochem. Commun. 9 (2007) 1486–1490, http://dx.doi.org/10.1016/j.elecom.2007. 02.008. A. Jain, S.P. Ong, G. Hautier, W. Chen, W.D. Richards, S. Dacek, S. Cholia, D. Gunter, D. Skinner, G. Ceder, K.A. Persson, Commentary: the Materials Project: a materials genome approach to accelerating materials innovation, Apl. Mater. 1 (2013) 11002, http://dx.doi.org/10.1063/1.4812323. A. Kuhn, O. Gerbig, C. Zhu, F. Falkenberg, J. Maier, B.V. Lotsch, Ultrafast Li electrolytes based on abundant Elements: Li10SnP2S12 and Li11Si2PS12, http://arxiv.org/abs/1402. 4586, (2014) 2-8. A. Kuhn, O. Gerbig, C. Zhu, F. Falkenberg, J. Maier, B.V. Lotsch, A new ultrafast superionic Li-conductor: ion dynamics in Li11Si2PS12 and comparison with other tetragonal LGPS-type electrolytes, Phys. Chem. Chem. Phys. 16 (2014) 14669–14674, http://dx.doi. org/10.1039/c4cp02046d. W.C. West, J.F. Whitacre, J.R. Lim, Chemical stability enhancement of lithium conducting solid electrolyte plates using sputtered LiPON thin films, J. Power Sources 126 (2004) 134–138, http://dx.doi.org/10.1016/j.jpowsour.2003.08.030. I. Tarhouchi, V. Viallet, P. Vinatier, M. Ménétrier, Electrochemical characterization of Li10SnP2S12: an electrolyte or a negative electrode for solid state Li-ion batteries? Solid State Ionics 296 (2016) 18–25, http://dx.doi.org/10.1016/j.ssi.2016.08.016. S. Hori, K. Suzuki, M. Hirayama, Y. Kato, T. Saito, M. Yonemura, R. Kanno, Synthesis, structure, and ionic conductivity of solid solution, Li10+δM1+δP2−δS12 (M = Si, Sn), Faraday Discuss 176 (2014) 83–94, http://dx.doi.org/10.1039/C4FD00143E. S.H. Kang, W.S. Yoon, K.W. Nam, X.Q. Yang, D.P. Abraham, Investigating the first-cycle irreversibility of lithium metal oxide cathodes for Li batteries, J. Mater. Sci. 43 (2008) 4701–4706, http://dx.doi.org/10.1007/s10853-007-2355-6. A. Sakuda, N. Nakamoto, H. Kitaura, A. Hayashi, K. Tadanaga, M. Tatsumisago, All-solidstate lithium secondary batteries with metal-sulfide-coated LiCoO2 prepared by thermal decomposition of dithiocarbamato complexes, J. Mater. Chem. 22 (2012) 15247, http:// dx.doi.org/10.1039/c2jm32490c. B. Chen, Z. Huang, X. Chen, Y. Zhao, Q. Xu, P. Long, S. Chen, X. Xu, A new composite solid electrolyte PEO/Li10GeP2S12/SN for all-solid-state lithium battery, Electrochim. Acta 210 (2016) 905–914, http://dx.doi.org/10.1016/j.electacta.2016.06.025. J. Li, L.E. Downie, L. Ma, W. Qiu, J.R. Dahn, Study of the failure mechanisms of LiNi0.8Mn0.1Co0.1O2 cathode material for lithium ion batteries, J. Electrochem. Soc. 162 (2015) A1401–A1408, http://dx.doi.org/10.1149/2.1011507jes. G. Govindasamy, P. Murugasen, S. Sagadevan, Optical and electrical properties of chemical bath deposited cobalt sulphide thin films, Mater. Res. 20 (2017) 62–67, http://dx. doi.org/10.1590/1980-5373-mr-2016-0441. S.S. Hegde, A.G. Kunjomana, K.A. Chandrasekharan, K. Ramesh, M. Prashantha, Optical and electrical properties of SnS semiconductor crystals grown by physical vapor deposition technique, Phys. B Condens. Matter 406 (2011) 1143–1148, http://dx.doi.org/10. 1016/j.physb.2010.12.068.
Contents lists available at ScienceDirect
Journal of Power Sources journal homepage: www.elsevier.com/locate/jpowsour
Electrochemical and interfacial behavior of all solid state batteries using Li10SnP2S12 solid electrolyte Carolina Vinadoa, Shanyu Wanga, Yang Hec, Xingcheng Xiaob, Yun Lia, Chongmin Wangc, Jihui Yanga,∗ a
Materials Science and Engineering Department, University of Washington, Seattle, WA, 98195, USA Chemical Sciences and Materials Systems Lab, General Motors R&D Center, Warren, MI, 48090, USA c Pacific Northwest National Laboratory, Environmental Molecular Science Laboratory, Richland, WA, 99354, USA b
H I GH L IG H T S
utilizing Li SnP S solid electrolyte in all-solid-state batteries. • Successfully inter-diffusion between LCO and LSPS causes a poor cycling stability. • Atomic • Li NbO coating improves the performance of the LCO/LSPS solid state batteries. 10
3
2 12
4
A R T I C LE I N FO
A B S T R A C T
Keywords: Li10SnP2S12 All-solid-state battery Solid electrolyte Atomic layer deposition Interfacial stability
Thio-Lithium Superionic Conductor (Thio-LISICON) Li10GeP2S12 equivalent Li10SnP2S12 (LSPS) is comparable in ionic conductivity yet with a lower cost as an electrolyte for all solid-state batteries (ASSBs). ASSBs with LSPS solid electrolyte (SE), lithium-indium alloy anode, and LiCoO2 (LCO) cathode were successfully fabricated and their electrochemical performance at 60 °C was examined. Atomic layer deposition of Li3NbO4 on LCO was conducted to improve the interfacial stability. The Li3NbO4 coating effectively improves the cycle stability of the ASSB. Electrochemical impedance spectroscopy tests indicate a rapid growth of charge transfer resistance upon cycling for the cell with the uncoated LCO, primarily due to the surface instability and build-up of a space charge layer. However, the ASSBs with Li3NbO4 coated LCO show a more stable interface with a negligible impedance increase upon cycling, attributable to the buffering and passivating roles of the Li3NbO4 coating. The interfacial microstructure was analyzed to elucidate at the underlying reasons for the impedance increase and the pivotal role of the Li3NbO4 coating.
1. Introduction Lithium ion batteries (LIBs) are widely commercialized to power portable devices. Due to their high energy density, it is the most attractive option for hybrid electric vehicles, electric vehicles, and other portable applications [1]. Given the sensitive and widespread nature of these applications, high energy LIBs with superior reliability and safety are desirable. Most commercially available LIBs use flammable liquid electrolytes, which risks fire or explosions in the event of a failure [2]. From a safety point of view, replacing the flammable liquid electrolyte with a solid one would be desirable. An eligible solid electrolyte (SE) should have an ionic conductivity comparable to that of current liquid electrolytes, e.g., 10−3-10−2 S cm−1 [2]. Other benefits of SEs include larger electrochemical windows and excellent thermal stability, as well
∗
as the possibility of high power and energy density, such as when paired with a lithium metal anode [3]. The recent finding of Li10GeP2S12 (LGPS), with an outstanding ionic conductivity of 12 mS cm−1, comparable to liquid electrolytes, has made the use of sulfide-based SEs an appealing option for all solid-state batteries (ASSBs) [4,5]. The high cost of Ge, however, largely limits the practical use of LGPS. Creating isostructural analogues by substituting Ge with Si, Al, or Sn, has been suggested [6–8]. The Sn analogue Li10SnP2S12 (LSPS) is the most promising of the group [8]. The main hurdle for developing successful ASSBs, however, is in minimizing the interfacial impedances between the SE and the electrodes, especially for the cathode/SE interfaces. The interfacial impedances primarily dominate the rate capability and cycling stability of the ASSBs [1,4,9]. First, SEs must be chemically and electrochemically
Corresponding author. E-mail address: [email protected] (J. Yang).
https://doi.org/10.1016/j.jpowsour.2018.06.038 Received 6 March 2018; Received in revised form 8 June 2018; Accepted 9 June 2018 0378-7753/ © 2018 Elsevier B.V. All rights reserved.
Please cite this article as: Vinado, C., Journal of Power Sources (2018), https://doi.org/10.1016/j.jpowsour.2018.06.038
Journal of Power Sources xxx (xxxx) xxx–xxx
C. Vinado et al.
Li-In foil as the anode. The working electrode and LSPS were placed in a stainless steel die with a diameter of 10 mm and pressed together under 500 MPa for 20 min. The cathode/LSPS pellet was then dissembled from the die and placed into an ASSC. 60 mg of In foil (0.62 V vs Li+/Li [27], 99.99%, CMR-Direct, USA) and 10 mg of Li foil (99.9%, MTI Corporation, USA) were placed on the top of the SE and a pressure of 20 MPa was then applied. All cell preparation processes were carried in a glovebox (Lab Star, Mbraun, Germany, H2O < 0.5 ppm, O2 < 0.5 ppm). The ASSCs under a 20 MPa pressure were then galvanostatically charged/discharged using a Neware cycler in the voltage range of 2.0–3.7 V vs. Li-In at 60 °C. AC EIS measurements were performed in a potentiostat (Versastat4, Ametek Scientific Instruments, USA). EIS data of a cold-pressed SE pellet were collected at room temperature in a symmetrical cell using carbon foil as the electrodes in the frequency range of 1–106 Hz and with an amplitude of 5 mV. Similarly, EIS tests of the ASSCs in the frequency range of 0.01–106 Hz and an amplitude of 10 mV were performed at various cycling stages. Cyclic voltammetry (CV) was carried out on LSPS pellets sandwiched between a stainless steel (SS) plate and Li or In foil as electrodes at a scanning rate of 10 mV s−1 from 0 to 5 V. The crystal structure of the materials was investigated by powder Xray diffraction (XRD, Bruker D8 Discover with IμS 2D detector, Germany). The LSPS powders were first placed in a capillary tube and sealed with wax before exiting the glovebox. A field emission SEM (FESEM, FEI Sirion XL30, USA) equipped with an Oxford EDS was used to image the particle morphology and obtain elemental analysis and mapping. The microstructures and compositions of the cycled electrodes were analyzed by a scanning TEM equipped with a high-angle annular dark-field (HAADF) detector and an EDS at the Pacific Northwest National Laboratory (STEM/HAADF/EDS, FEI Titan 80–300 kV, USA). In order to reduce air- or moisture-exposure and avoid the unintended contaminations, the cycled cells were disassembled, the cathodes were scraped, and the cathode powders were placed on a TEM mesh in an Ar filled glovebox. The TEM holder was sealed and rapidly transferred into the TEM chamber with a minimal exposure.
stable against the electrode materials and able to form a stable and conductive interface (solid electrolyte interphase-like interface layer) [9–13]. Based on the theoretical calculations, sulfide-based SEs show narrow electrochemical windows, such as ∼1.8–2.3 V vs. Li+/Li for LGPS [14]. The narrow electrochemical windows of sulfides have been experimentally verified, although these had been largely overestimated by the conventional experimental method of using a Li/electrolyte/ inert metal semi-blocking electrode with limited contact areas [15]. In addition, due to the difference in Li+ chemical potential and resultant asymmetric Li+ migration, the depletion of Li+ on the sulfide side and the formation of a space charge layer largely increase the interfacial resistance [16]. To facilitate successful ASSBs, interfacial modifications on either electrodes or SE are necessary to mitigate the chemical diffusion and space charge layer, thus reducing the interfacial impedance and extending the electrochemical window beyond 4.2 V (vs. Li+/Li) [9–13,17–20]. To mitigate the interfacial resistance, the coating material should be a good ionic conductor, while thin and uniform. Typically, wet chemical methods such as sol-gel or spray coating, are used for creating surface coatings; these methods, however, leave non-uniform and thick coatings [21]. Alternatively, atomic layer deposition (ALD) boasts excellent conformity, atomic scale thickness control, and uniformity over large areas, emerging as a promising process for battery electrode coating [9,22]. The sequential and self-limiting nature of ALD allows for ultra-thin and pinhole free coatings to be applied. These subnanometer coating layers have the added benefits of lessened mechanical stresses due to the crystal mismatch and reduced thickness constraints as far as the electron and lithium ion conductivities of the coating material are concerned [21]. To evaluate the potential of LSPS as an SE and study its interfacial behavior with typical cathode materials, ASSBs with LSPS SE, Li/In anode, pristine and ALD Li3NbO4-coated LiCoO2 cathodes were fabricated. Li3NbO4 was chosen because of its reasonable ionic conductivity and excellent chemical stability against both layered oxide cathodes and LSPS, thus effectively protecting the cathode from undesirable side reactions with the electrolytes [23–26]. The electrochemical performance and interfacial characteristics of the ASSBs using pristine and coated LiCoO2 were systematically explored by galvanostatic cycling, AC electrochemical impedance spectroscopy (EIS), scanning electron microscope (SEM), transmission electron microscope (TEM), energy dispersive X-ray spectroscopy (EDS), etc. The study shows that the ALD coating largely improves the electrochemical performance of ASSBs with the LSPS SE, in large part due to the improved interfacial compatibility and decreased charge transfer impedance.
3. Results and discussion 3.1. Characterization of the solid electrolyte and cathode materials Fig. 1a shows the crystal structure of LSPS, taken from the online database Materials Project [28]. LSPS crystalizes in a tetragonal structure with a space group P42/mc. The lattice parameters are a = 8.854 Å and c = 12.851 Å. The three-dimensional percolating structure ensures high Li+ conductivity. To verify the crystal structure of the as-received LSPS, powder XRD was carried out, as shown in Fig. 1b. Based on the above crystal structure, the standard XRD pattern of LSPS was also simulated using the Materials Studio's software, also shown in Fig. 1b. Our simulated XRD matches well with the result presented in Materials Project and those given by Bron, Kuhn, and Tarhouchi et al. [7,29,30]. As shown, the experimental XRD can be well indexed to the simulated XRD, indicating good phase purity of the LSPS powders. In addition, the sharp and well-defined diffraction peaks indicate good crystallinity of the SE. Fig. 1c shows the Nyquist plot for an LSPS pellet cold pressed at 500 MPa. Similar to the previous work done by Bron et al. [7], we observed grain resistance and relatively small grain boundary contribution, even without heat treatment. The Nyquist plot can be well fitted by the equivalent circuit shown in the inset of Fig. 1b. Rb, corresponding to the high-frequency intercept with the real axis, is the bulk resistance; the inductance segment L is mainly due to the cabling and leads of the set up [31]; the Rgb|CPEgb combination, showing as the semicircle in the mid-low frequency range, corresponds to the ionic resistance and capacitance of the grain boundaries in the LSPS pellet [8]. Based on the bulk resistance and geometry of the pellet, a total
2. Experimental ALD of ultrathin lithium niobium oxide coatings were deposited using a Savannah 100 Atomic Layer Deposition system (Cambridge Nanotech Inc., USA), with lithium tert-butoxide (LiOtBu, (CH3)3COLi, Sigma Aldrich, USA) and niobium ethoxide [C10H25NbO5, Sigma Aldrich] as the precursors. The distilled water was used as the oxidant and Argon (99.99%) as the carrier gas. The pulse, exposure, and purge times for all precursors are 0.1 s, 10 s, and 20 s, respectively, in each sub-cycle. The heating temperature was 160 °C for both Li and Nb precursors, and 80 °C for water. The deposition temperature was around 200 °C for the LiCoO2 powder samples (LCO, 99.8%, Sigma Aldrich, USA). The electrochemical properties LCO/LSPS/Li-In ASSBs were investigated using a homemade all-solid-state cells shown in Fig. S1 (Supporting Information, SI). Electrode mixtures consisted of pristine LCO or Li3NbO4 coated LCO (c-LCO) powders, and LSPS powders (Nanomyte SSE-10, NEI Corporation, USA) in a weight ratio of 70:30. They were prepared by hand grinding in a mortar for 30 min. Twoelectrode all solid state cells (ASSCs) were fabricated using ∼ 12.7 mg of electrode mixture as the cathode, 80 mg LSPS as the separator, and 2
Journal of Power Sources xxx (xxxx) xxx–xxx
C. Vinado et al.
Fig. 1. (A) Crystal structure of LSPS. [SnS4] (Orange) and [PS4] (Aqua) tetrahedra are shown. (b) Experimental and simulated powder X-ray diffraction (XRD) of LSPS. (c) Nyquist plot of the experimental (Black Square) and fit (Red line) of a LSPS pellet in a C/LSPS/C cell at room temperature. (d) Cyclic voltammograms of the Li/LSPS/SS and In/ LSPS/SS for the electrochemical stability window evaluation of the SE vs. Li & In between −0.5 and 5 V vs. Li+/Li at room temperature. Scanning electron microscopy images of the (e) LSPS, (f) uncoated LCO, and (g) c-LCO powders. (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.)
ionic conductivity of ∼3 mS cm−1 at room temperature can be obtained for our cold-pressed LSPS sample, comparable to the literature values [7,8,30,32,33]. To select an appropriate anode material, cyclic voltammetry was done on a cold pressed LSPS pellet. Fig. 1d shows the cyclic voltammograms of LSPS to evaluate the electrochemical stability window in Li/LSPS/SS and In/LSPS/SS asymmetric cells. The red line is the CV curve for the Li/LSPS/SS setup. A large cathodic peak, due to the deposition of metallic lithium (Li+ + e−→ Li), and a slight anodic peak associated with the stripping of metallic lithium (Li → Li+ + e−), are observed at around 0 V, similar to those of related systems [30,32,33]. On subsequent cycles, the peak's currents gradually decrease (see details: Fig. S2). It is well known that sulfide-based SEs are unstable against Li metal, which can cause a reduction of the SE upon contact with Li metal [30]. This is verified by the significant loss of current signal on the following cycles, further ascribed to the largely increased resistance from the interfacial reaction products. For the In/LSPS/SS setup, however, no noticeable change in the cathodic and anodic peaks can be observed during the voltage sweep (black line in Fig. 1d), implying no redox occurring for the In/LSPS interface. Thus, Li-In alloys with a high oxidation potential of ∼0.62 V vs. Li+/Li and good ductility were chosen for our experiments to ensure good contact with LSPS and bypass any confounding factors arising from parasitic reactions during our electrochemical tests. Fig. 1e shows a SEM image of the LSPS SE as received. The particles are irregularly shaped oblongs with sizes of 1–10 μm and significant aggregations. EDS (Fig. S3) shows a fairly uniform elemental distribution, consistent with the single phase XRD pattern. Similarly, Fig. 1f and g shows the SEM images of the pristine and Li3NbO4-coated LCO, respectively, both of which have a chunk or irregular-shape morphology
and a wide size distribution of 1–20 μm. EDS mapping of the coated sample shows a fairly even elemental distribution (Fig S4 in SI), consistent with its single phase XRD (Fig.S5 in SI). The coating layer is undetectable under SEM/EDS or XRD, primarily due to the thin nature and low volume percentage content of the coating (only 5 ALD cycles and ∼0.5 nm thick). 3.2. Electrochemical performance of the all-solid-state cells In order to study the electrochemical performance of the SE and its compatibility with LCO, ASSBs with pristine and c-LCO were fabricated and tested. Fig. 2a shows a comparison of the cycling stability between LCO and c-LCO. First, LCO shows a low discharge capacity of ∼105 mA h/g at C/10 and 60 °C, much lower than the value achieved in liquid cells (∼145 mA h/g at an upper cut-off of 4.3 V vs. Li+/Li) [32]. More importantly, the initial Coulombic efficiency is only 82%, also much lower than that of liquid cells (98%) [34]; similarly, the system in general possesses other limitations when compared to liquid electrolyte systems, such as added interfacial resistances, impaired kinetics under certain circumstances, and a thicker electrolyte layer. The significant initial capacity loss indicates an interfacial instability between LSPS and LCO that consumes Li in order to form a solid electrolyte interphase (SEI). The large voltage polarization of the LCO cell also indicates its large interfacial impedance. The LCO cell activates its discharge capacity to ∼110 mA h/g in the 2nd cycle and then experiences a severe capacity decay, only retaining ∼60% initial capacity after 70 cycles. This large capacity decay is accompanied by a large voltage decline. For instance, the average discharge voltage increases from 3.24 V for the 1st cycle to 3.28 V on the 10th cycle and gradually decreases to 3.15 V by the 50th cycle., as seen in Fig. 2b. The large 3
Journal of Power Sources xxx (xxxx) xxx–xxx
C. Vinado et al.
Fig. 2. (A) Cycling performance of LCO/LSPS/Li-In and c-LCO/LSPS/Li-In between 2.0 and 3.7 V vs. Li-In (2.62–4.32 V vs Li+/Li) at C/10 (0.13 mA/cm2) and 60 °C. Voltage profiles of (b) LCO/LSPS/Li-In and (c) c-LCO/LSPS/Li-In.
interdiffusion are allowed to occur for a longer time when using a low C-rate, resulting in a severally deteriorated performance. The cell shows even lower capacities at higher C-rates: ∼20 mA h/g at C/10, 10 mA h/ g at C/5, and no observable capacity at C/3 or higher. In addition, the LCO cell only retains 15 mA h/g discharge capacity when cycled at C/ 10 again, demonstrating a failure of the cell. Conversely, the c-LCO cell has a much better rate capability, showing a discharge capacity of ∼120 mA h/g at C/20, 115 mA h/g at C/10, 95 mA h/g at C/5, 70 mA h/g at C/3, and 20 mA h/g at 1C, as shown in Fig. 3a. After the high-rate cycles, the c-LCO cell can recover ∼95% initial capacity at C/ 10, corroborating its much better cyclability. A comparison of voltage profiles, shown in Fig. 3b, clearly indicate higher discharge voltages and lower overpotentials of the c-LCO cell, as compared with the LCO cell. The noticeably improved cycling stability and rate capability of the c-LCO cell are primarily originated from the reduced interfacial resistance and improved interfacial stability, which will be discussed in more details below.
voltage decline is a consequence of a radically increasing LCO/LSPS interfacial resistance, contributing to the severe capacity decay, together with the active material loss, as LCO is consumed by the reactions with LSPS, which will be discussed later. The c-LCO cell not only shows a higher discharge specific capacity but also better cycling stability. The initial specific discharge capacity and Coulombic efficiency of the c-LCO cell are ∼118 mA h/g and ∼87%, respectively. Its capacity gradually activates to ∼127 mA h/g by the 5th cycle, consistent with the decreasing overpotential and interfacial impedance of the cell. The activation process is presumably originated from the interface activation (improved particle contact, SEI formation, void filling, etc.). Since Li3NbO4 is stable against both LCO and LSPS, the coating largely avoids the direct physical contact of LCO and LSPS and thus mitigates their chemical reactions, protecting the interface in the following cycles. In addition, due to the buffering effect of the coating, the space charge layer and resultant resistance are largely reduced [16]. These two effects contribute to stable LCO/Li3NbO4/ LSPS interfaces and thus a much better cycling stability. The c-LCO cell can retain ∼85% initial capacity after 70 C/10 cycles, outperforming the LCO cell. More importantly, the average discharge voltage of the cLCO cell initially increases from 3.14 V for the 1st cycle to 3.42 V for the 10th cycle, and slowly declines to 3.24 V for the 50th cycle, as shown on Fig. 2c. The initial capacities of the uncoated and coated LCO, as well as their cycling behavior is well consistent with those of other LCO ASSBs [12]. Fig. 3a shows the rate capability of the LCO and c-LCO ASSBs. Due to severe interfacial instability and thus high resistance, pristine LCO shows poor rate capability and cycling stability, as shown in Fig. 3a. The LCO cell only shows a low initial capacity of ∼77 mA h/g and loses almost 75% of its capacity in first 5 cycles at C/20. The extremely low initial capacity may be related to severe side reactions of LCO/LSPS at the slow charge process (C/20), as the chemical reactions and elemental
3.3. Electrochemical impedance spectroscopy study As demonstrated above, surface coating to physically separate LCO and LSPS is indispensable for employing this highly conductive and technically feasible SE. The improved interfacial stability is evidenced by the enhanced cycling stability and rate capability, which can be further clarified by monitoring the impedance evolution with cycling. Fig. 4a and b depict the room temperature Nyquist plots of the LCO and c-LCO cells at different cycling stages. The Nyquist plots could be well fitted by the equivalent circuit shown in the inset of Fig. 4a, which contains the serial combination of internal resistance RS (high frequency intercept with real axis), interfacial resistance Rsei (high frequency semicircle, ∼100 kHz), charge transfer resistance Rct (low frequency semicircle, ∼500–600 Hz), and Warburg element representing 4
Journal of Power Sources xxx (xxxx) xxx–xxx
C. Vinado et al.
Fig. 3. (A) Rate capability of LCO/LSPS/Li-In and c-LCO/LSPS/Li-In between 2.0 and 3.7 V vs. Li-In (2.62–4.32 V vs. Li+/Li) at 60 °C. (b) Voltage profiles at different rates of LCO/LSPS/Li-In (dashed lines) and c-LCO/LSPS/Li-In (solid lines).
indicates the interfacial instability between LCO and LSPS, due to chemical reactions, mutual elemental diffusions, etc. The interfacial analyses will be presented below to clarify this. The thin Li3NbO4 coating does improve the interfacial stability of LCO/LSPS, although it cannot completely get rid of the Rsei. Similar effects can be found for ASSBs with other coated materials [9–13,18,20]. Further modification of the coating thickness is required to mitigate the SEI formation, and it is in our plans for future work. Large charge transfer resistances Rct, ∼13700 Ω cm and 7000 Ω cm for the LCO and c-LCO cells, respectively, can be observed which can be attributed to the resistive surface layers originated from surface contamination of LSPS and the reactivity of LCO/LSPS. The resistive surface layers not only impede Li+ transfer but also limit the electron transfer across active particle surfaces. Similar to the Rsei, Rct also experiences an activation process during the first several cycles; upon assembly of the cell, a chemically inactive surface layer forms; once the cell is cycled, this layer is destroyed and modified such that it is less resistive and facilitates the charge transfer process [26]. This is referred as the ‘activation process’, and the reason for the initial decrease in impedance. After the fifth cycle, Rct reduces noticeably to 1700 and 900 Ω cm for the LCO and c-LCO cells, respectively. The decrease in Rct should share the same origin as that of Rsei, associated with the removal/formation of the passivation layer on the particle surfaces and/ or increased contact area. In addition, the lower Rct of the c-LCO cell is mainly ascribed to the surface protection by the coating, which largely
the solid-state ion diffusion (low frequency tail), as well as the respective constant phase elements (CPE) associated with Rsei and Rct [17–20]. The fitting results of RS, Rsei, and Rct are summarized in Fig. 4c. The internal resistance of the as-prepared cells, including the resistances from cathode, anode, separator, current collectors, and leads, are ∼560 Ω cm for the LCO cell and ∼790 Ω cm for the c-LCO cell; the higher RS of the c-LCO cell is mainly due to the electronically insulating coating. The significant RS drop after the 1st cycle for both cells is primarily attributed to improved contact and the activation of particles after high-temperature cycling. The c-LCO cell keeps its low RS (210 Ω cm) after 10 cycles, while the LCO cell shows largely increased RS (680 Ω cm), consistent with its interfacial instability. Due to the air- and moisture-sensitivity of LSPS, the LSPS particle surface is inevitably contaminated and forms a passivating layer. Although LCO is stable in air, its surface could also absorb trace amounts of H2O, which further contaminates the LSPS surface when mixing and assembling them together. These facts could account for the large Rsei of the as-assembled LCO (∼2500 Ω cm) and c-LCO cells (∼3530 Ω cm). Interestingly, after one charge at 60 °C Rsei of both cells almost disappears (∼200 Ω cm). The exact origin of this phenomenon, however, needs further experimental and theoretical efforts. Rsei increases slightly to ∼450 Ω cm after 5 cycles, however, it jumps to large values of ∼6740 Ω cm and ∼2590 Ω cm for the LCO and c-LCO cells after 10 cycles, respectively. The significant increase in Rsei clearly
Fig. 4. Room Temperature Nyquist plots of (a) LCO/LSPS/Li-In and (b) c-LCO/LSPS/Li-In at different cycling stages (as-prepared and after 1st, 5th, and 10th charge). (c) Summary of Rs, Rsei, and Rct with cycling derived from EIS fitting. 5
Journal of Power Sources xxx (xxxx) xxx–xxx
C. Vinado et al.
electron transport [37,38], resulting in the high resistance and considerable overpotential of the LCO/LSPS/Li-In cell. Fig. 5b shows a TEM image of the c-LCO/LSPS interface after 10 C/ 10 cycles at 60 °C. Due to the ultrathin coating and beam-sensitivity of Li-ion conductors, Li3NbO4 is undetectable. A clear and narrow interface, however, can be observed, as compared with that of the LCO/LSPS interface. The EDX line scan of the Co, Sn, and P elements across the interface shows a thin inter-diffusion length of ∼4 nm. Co diffusion into LSPS was mostly blocked, and P and Sn diffuse ∼4 nm into c-LCO. It is obvious that the ultrathin coating layer largely impedes the atomic inter-diffusion, resulting in much thinner “damage layer” and hence lower interfacial resistance compared with uncoated particles. The Li3NbO4 coating, as suggested [39], also presumably alleviates the space charge layer formation, further reducing the Li+ transfer resistance. As shown in Fig. 4c, the c-LCO cell still suffers large impedance growth after 10 cycles, probably due to the ultrathin coating in this work. Further optimization of the coating thickness is required to optimize the interface behavior and cycling performance. Based on the above results, in order to use the highly conductive and cost-effective LSPS in ASSBs, surface protection of the cathode materials is necessary, at least for the layered cathodes, i.e., LCO, NMC (LiNi1-x-yMnxCoyO2), NCA (LiNi0.8Co0.15Al0.05O2). The high reactivity of sulfide-based SEs mainly originates from the high malleability of S2−, which also gives rise to its merits: high ionic conductivity and good deformability, as compared with oxide-based SEs. Here we have successfully demonstrated that an ALD coating with ultrathin oxide-based Li-ion conductors is effective at not only protecting the active-material/ SE interface, but also reducing interfacial resistance, thus mitigating the cell impedance and its build-up with cycling.
prevents the formation of a thick SEI [3,9,32]. The activation process, including the reduction of all three resistances, agrees well with the reduction of overpotential and increased specific capacity in the initial several cycles, as shown in Fig. 2. The capacity and impedance activation process is widely observed in ASSBs [35] and liquid cells with surface sensitive active materials, such as NMC811 (LiNi0.8Mn0.1Co0.1O2) [36]. After 10 cycles, Rct increases significantly to ∼9800 Ω cm and 4200 Ω cm for the LCO and c-LCO cells, respectively. It is obvious that the SEI and charge transfer resistances dominate the cell performances, including the cycling stability and rate capability. After 10 cycles, the LCO cell shows a large total resistance of ∼17000 Ω cm, much larger than that of the c-LCO cell (∼7000 Ω cm). The resistance of the LCO cell will keep increasing without forming stable interfaces, and finally terminate the cell due to a strong voltage polarization. Surface coating with Li3NbO4 largely improves the interfacial stability, and thus the cycling stability and rate capability, as demonstrated in other studies [3,9,13,17,34].
3.4. TEM analysis of LCO/LSPS and c-LCO/LSPS interfaces To gain better understanding of the LCO/LSPS interfacial behavior and the effects of the surface coating, TEM observations of the composite cathodes were carried out to directly visualize the LCO/LSPS interface. The ASSBs were disassembled after 10 C/10 cycles at 60 °C in an argon-filled glove box. Figure 5a shows a TEM image of an LCO/ LSPS interface, and a thick passivating layer with gradually varying contrast can be observed in between. An EDX line scan across the interface confirms the interface thickness to be of ∼18 nm. The variations of Co, Sn, and P elements at the interface show that the mutual diffusion length is ∼5 nm on both sides. Particularly, Sn is largely depleted in the LSPS side, indicating that Sn deeply migrates into LCO. A high-resolution TEM image, shown in Fig. S6 (SI), indicates that the interface seems to be composed of crystalline nanoparticles, which are most likely cobalt and/or tin sulfides. Similar results have been found for Li2S-P2S5 SEs interphases [13,18]. The sulfides, which presumably make up this interfacial layer, are highly resistive for both Li+ and
4. Conclusions All solid-state batteries with LiCoO2 cathodes, Li10SnP2S12 solid electrolyte, and Li-In anode were fabricated to evaluate the electrochemical performance of low-cost Li10SnP2S12 and its interfacial behavior with LiCoO2. Poor cycling stability and rate capability can be
Fig. 5. TEM images and corresponding EDX line scans of (a) LCO/LSPS and (b) c-LCO/LSPS interfaces after 10 C/10 cycles at 60 °C. 6
Journal of Power Sources xxx (xxxx) xxx–xxx
C. Vinado et al.
observed for the cell with pristine LiCoO2, primarily due to the instability of the LCO/LSPS interface. Considerable atomic inter-diffusion causes a thick interphase layer and markedly increases the interfacial resistance, resulting in large voltage polarizations and poor cyclability. Surface coating of ultrathin Li3NbO4 on LiCoO2 by atomic layer deposition largely improves the cycling and rate performance of the ASSBs. The enhanced electrochemical performance is further traced back to the decreased charge transfer and interfacial resistances, originated from the mitigation of atomic inter-diffusion and the formation of resistive interphases. Our study also demonstrated that ALD coating with thin layer oxide-based Li-ion conductors is imperative and effective for the utilization of sulfide-based SEs for ASSBs.
[14]
[15] [16]
[17]
[18]
[19]
Acknowledgements [20]
CMW thanks the support of the Assistant Secretary for Energy Efficiency and Renewable Energy, Office of Vehicle Technologies of the U.S. Department of Energy under Contract No. DE-AC02-05CH11231, Subcontract No. 18769 and No. 6951379 under the Advanced Battery Materials Research (BMR) program. The microscopic analysis in this work was conducted in the William R. Wiley Environmental Molecular Sciences Laboratory (EMSL), a national scientific user facility sponsored by DOE's Office of Biological and Environmental Research and located at PNNL. PNNL is operated by Battelle for the Department of Energy under Contract DE-AC05-76RLO1830. X.X also acknowledges the support by the Assistant Secretary for Energy Efficiency and Renewable Energy, Vehicle Technologies Office of the U.S. Department of Energy under Contract No. Award Number DE-EE0007787 under the Battery Material Research (BMR) Program.
[21]
[22]
[23]
[24]
[25]
[26]
Appendix A. Supplementary data Supplementary data related to this article can be found at http://dx. doi.org/10.1016/j.jpowsour.2018.06.038.
[27]
References
[28]
[1] B. Wu, S. Wang, W.J. Evans IV, D.Z. Deng, J. Yang, J. Xiao, Interfacial behaviours between lithium ion conductors and electrode materials in various battery systems, J. Mater. Chem. 4 (2016) 15266–15280, http://dx.doi.org/10.1039/C6TA05439K. [2] J.B. Goodenough, Y. Kim, Challenges for rechargeable Li batteries, Chem. Mater. 22 (2010) 587–603, http://dx.doi.org/10.1021/cm901452z. [3] A. Sakuda, A. Hayashi, T. Ohtomo, S. Hama, M. Tatsumisago, All-solid-state lithium secondary batteries using LiCoO2 particles with pulsed laser deposition coatings of Li2S–P2S5 solid electrolytes, J. Power Sources 196 (2011) 6735–6741, http://dx.doi.org/ 10.1016/j.jpowsour.2010.10.103. [4] N. Kamaya, K. Homma, Y. Yamakawa, M. Hirayama, R. Kanno, M. Yonemura, T. Kamiyama, Y. Kato, S. Hama, K. Kawamoto, A. Mitsui, A lithium superionic conductor, Nat. Mater. 10 (2011) 682–686, http://dx.doi.org/10.1038/nmat3066. [5] Y. Kato, S. Hori, T. Saito, K. Suzuki, M. Hirayama, A. Mitsui, M. Yonemura, H. Iba, R. Kanno, High-power all-solid-state batteries using sulfide superionic conductors, Nat. Energy (2016) 16030, http://dx.doi.org/10.1038/nenergy.2016.30. [6] S.P. Ong, Y. Mo, W.D. Richards, L. Miara, H.S. Lee, G. Ceder, Phase stability, electrochemical stability and ionic conductivity of the Li10 ± 1MP2X12 (M = Ge, Si, Sn, Al or P, and X = O, S or Se) family of superionic conductors, Energy Environ. Sci. 6 (2013) 148–156, http://dx.doi.org/10.1039/C2EE23355J. [7] P. Bron, S. Johansson, K. Zick, J. Schmedt auf der Günne, S. Dehnen, B. Roling, Li10SnP2S12: an affordable lithium superionic conductor, J. Am. Chem. Soc. 135 (2013) 15694–15697, http://dx.doi.org/10.1021/ja407393y. [8] P. Bron, S. Dehnen, B. Roling, Li10Si0.3Sn0.7P2S12 – a low-cost and low-grain-boundaryresistance lithium superionic conductor, J. Power Sources 329 (2016) 530–535, http:// dx.doi.org/10.1016/j.jpowsour.2016.08.115. [9] J.H. Woo, J.E. Trevey, A.S. Cavanagh, Y.S. Choi, S.C. Kim, S.M. George, K.H. Oh, S. Lee, Nanoscale Interface Modification of LiCoO2 by Al2O3 Atomic Layer Deposition for Solidstate Li Batteries, vol. 159, (2012), pp. 1120–1124, http://dx.doi.org/10.1149/2. 085207jes. [10] N. Ohta, K. Takada, L. Zhang, R. Ma, M. Osada, T. Sasaki, Enhancement of the high-rate capability of solid-state lithium batteries by nanoscale interfacial modification, Adv. Mater. 18 (2006) 2226–2229, http://dx.doi.org/10.1002/adma.200502604. [11] K. Takada, T. Ohno, N. Ohta, T. Ohnishi, Y. Tanaka, Positive and negative aspects of interfaces in solid-state batteries, ACS Energy Lett. 3 (2018) 98–103, http://dx.doi.org/ 10.1021/acsenergylett.7b01105. [12] A. Sakuda, H. Kitaura, A. Hayashi, K. Tadanaga, M. Tatsumisago, J. Power Sources 189 (2009) 527–530. [13] A. Sakuda, H. Kitaura, A. Hayashi, K. Tadanaga, M. Tatsumisago, All-solid-state lithium
[29]
[30]
[31]
[32]
[33]
[34]
[35]
[36]
[37]
[38]
[39]
7
secondary batteries with oxide-coated LiCoO2 electrode and Li2S–P2S5 electrolyte, J. Power Sources 189 (2009) 527–530, http://dx.doi.org/10.1016/j.jpowsour.2008.10.129. Y. Zhu, X. He, Y. Mo, First principles study on electrochemical and chemical stability of solid electrolyte–electrode interfaces in all-solid-state Li-ion batteries, J. Mater. Chem. 4 (2016) 3253–3266, http://dx.doi.org/10.1039/C5TA08574H. F. Han, T. Gao, Y. Zhu, K.J. Gaskell, C. Wang, A battery made from a single material, Adv. Mater. 27 (2015) 3473–3483, http://dx.doi.org/10.1002/adma.201500180. J. Haruyama, K. Sodeyama, L. Han, K. Takada, Y. Tateyama, Space–charge layer effect at interface between oxide cathode and sulfide electrolyte in all-solid-state lithium-ion battery, Chem. Mater. 26 (2014) 4248–4255, http://dx.doi.org/10.1021/cm5016959. Y. Iriyama, H. Kurita, I. Yamada, T. Abe, Z. Ogumi, Effects of surface modification by MgO on interfacial reactions of lithium cobalt oxide thin film electrode, J. Power Sources 137 (2004) 111–116, http://dx.doi.org/10.1016/j.jpowsour.2004.05.029. Y. Huang, J. Chen, J. Ni, H. Zhou, X. Zhang, A modified ZrO2-coating process to improve electrochemical performance of Li(Ni1/3Co1/3Mn1/3)O2, J. Power Sources 188 (2009) 538–545, http://dx.doi.org/10.1016/j.jpowsour.2008.12.037. Y. Huang, J. Chen, F. Cheng, W. Wan, W. Liu, H. Zhou, X. Zhang, A modified Al2O3 coating process to enhance the electrochemical performance of Li(Ni1/3Co1/3Mn1/3)O2 and its comparison with traditional Al2O3 coating process, J. Power Sources 195 (2010) 8267–8274, http://dx.doi.org/10.1016/j.jpowsour.2010.07.021. J.-H. Shim, K.-S. Lee, A. Missyul, J. Lee, B. Linn, E.C. Lee, S. Lee, Characterization of spinel LixCo2O4 -coated LiCoO2 prepared with post-thermal treatment as a cathode material for lithium ion batteries, Chem. Mater. 27 (2015) 3273–3279, http://dx.doi.org/10. 1021/acs.chemmater.5b00159. Z. Chen, Y. Qin, K. Amine, Y.-K. Sun, Role of surface coating on cathode materials for lithium-ion batteries, J. Mater. Chem. 20 (2010) 7606, http://dx.doi.org/10.1039/ c0jm00154f. H. Kaneda, Y. Koshika, T. Nakamura, H. Nagata, R. Ushio, K. Mori, Improving the cycling performance and thermal stability of LiNi0.6Co0.2Mn0.2O2 cathode materials by Nb-doping and surface modification, Int. J. Electrochem. Sci. 12 (2017) 4640–4653, http://dx.doi. org/10.20964/2017.06.19. K. Takada, N. Ohta, L. Zhang, K. Fukuda, I. Sakaguchi, R. Ma, M. Osada, T. Sasaki, Interfacial modification for high-power solid-state lithium batteries, Solid State Ionics 179 (2008) 1333–1337, http://dx.doi.org/10.1016/j.ssi.2008.02.017. B. Ahmed, C. Xia, H.N. Alshareef, Electrode surface engineering by atomic layer deposition: a promising pathway toward better energy storage, Nano Today 11 (2016) 250–271, http://dx.doi.org/10.1016/j.nantod.2016.04.004. K. Takada, N. Aotani, K. Iwamoto, S. Kondo, Solid state lithium battery with oxysulfide glass, Solid State Ionics 86 (1996) 877–882, http://dx.doi.org/10.1088/1751-8113/44/ 8/085201. S. Liu, X. Chen, J. Zhao, J. Su, C. Zhang, T. Huang, J. Wu, A. Yu, Uncovering the role of Nb modification in improving the structure stability and electrochemical performance of LiNi0.6Co0.2Mn0.2O2 cathode charged at higher voltage of 4.5 V, J. Power Sources 374 (2018) 149–157, http://dx.doi.org/10.1016/j.jpowsour.2017.11.037. N. Ohta, K. Takada, I. Sakaguchi, L. Zhang, R. Ma, K. Fukuda, M. Osada, T. Sasaki, LiNbO3-coated LiCoO2 as cathode material for all solid-state lithium secondary batteries, Electrochem. Commun. 9 (2007) 1486–1490, http://dx.doi.org/10.1016/j.elecom.2007. 02.008. A. Jain, S.P. Ong, G. Hautier, W. Chen, W.D. Richards, S. Dacek, S. Cholia, D. Gunter, D. Skinner, G. Ceder, K.A. Persson, Commentary: the Materials Project: a materials genome approach to accelerating materials innovation, Apl. Mater. 1 (2013) 11002, http://dx.doi.org/10.1063/1.4812323. A. Kuhn, O. Gerbig, C. Zhu, F. Falkenberg, J. Maier, B.V. Lotsch, Ultrafast Li electrolytes based on abundant Elements: Li10SnP2S12 and Li11Si2PS12, http://arxiv.org/abs/1402. 4586, (2014) 2-8. A. Kuhn, O. Gerbig, C. Zhu, F. Falkenberg, J. Maier, B.V. Lotsch, A new ultrafast superionic Li-conductor: ion dynamics in Li11Si2PS12 and comparison with other tetragonal LGPS-type electrolytes, Phys. Chem. Chem. Phys. 16 (2014) 14669–14674, http://dx.doi. org/10.1039/c4cp02046d. W.C. West, J.F. Whitacre, J.R. Lim, Chemical stability enhancement of lithium conducting solid electrolyte plates using sputtered LiPON thin films, J. Power Sources 126 (2004) 134–138, http://dx.doi.org/10.1016/j.jpowsour.2003.08.030. I. Tarhouchi, V. Viallet, P. Vinatier, M. Ménétrier, Electrochemical characterization of Li10SnP2S12: an electrolyte or a negative electrode for solid state Li-ion batteries? Solid State Ionics 296 (2016) 18–25, http://dx.doi.org/10.1016/j.ssi.2016.08.016. S. Hori, K. Suzuki, M. Hirayama, Y. Kato, T. Saito, M. Yonemura, R. Kanno, Synthesis, structure, and ionic conductivity of solid solution, Li10+δM1+δP2−δS12 (M = Si, Sn), Faraday Discuss 176 (2014) 83–94, http://dx.doi.org/10.1039/C4FD00143E. S.H. Kang, W.S. Yoon, K.W. Nam, X.Q. Yang, D.P. Abraham, Investigating the first-cycle irreversibility of lithium metal oxide cathodes for Li batteries, J. Mater. Sci. 43 (2008) 4701–4706, http://dx.doi.org/10.1007/s10853-007-2355-6. A. Sakuda, N. Nakamoto, H. Kitaura, A. Hayashi, K. Tadanaga, M. Tatsumisago, All-solidstate lithium secondary batteries with metal-sulfide-coated LiCoO2 prepared by thermal decomposition of dithiocarbamato complexes, J. Mater. Chem. 22 (2012) 15247, http:// dx.doi.org/10.1039/c2jm32490c. B. Chen, Z. Huang, X. Chen, Y. Zhao, Q. Xu, P. Long, S. Chen, X. Xu, A new composite solid electrolyte PEO/Li10GeP2S12/SN for all-solid-state lithium battery, Electrochim. Acta 210 (2016) 905–914, http://dx.doi.org/10.1016/j.electacta.2016.06.025. J. Li, L.E. Downie, L. Ma, W. Qiu, J.R. Dahn, Study of the failure mechanisms of LiNi0.8Mn0.1Co0.1O2 cathode material for lithium ion batteries, J. Electrochem. Soc. 162 (2015) A1401–A1408, http://dx.doi.org/10.1149/2.1011507jes. G. Govindasamy, P. Murugasen, S. Sagadevan, Optical and electrical properties of chemical bath deposited cobalt sulphide thin films, Mater. Res. 20 (2017) 62–67, http://dx. doi.org/10.1590/1980-5373-mr-2016-0441. S.S. Hegde, A.G. Kunjomana, K.A. Chandrasekharan, K. Ramesh, M. Prashantha, Optical and electrical properties of SnS semiconductor crystals grown by physical vapor deposition technique, Phys. B Condens. Matter 406 (2011) 1143–1148, http://dx.doi.org/10. 1016/j.physb.2010.12.068.