Capacity Fade in Solid-State Batteries: Interphase Formation and

Article pubs.acs.org/cm

Capacity Fade in Solid-State Batteries: Interphase Formation and Chemomechanical Processes in Nickel-Rich Layered Oxide Cathodes and Lithium Thiophosphate Solid Electrolytes Raimund Koerver,† Isabel Aygün,† Thomas Leichtweiß,‡ Christian Dietrich,† Wenbo Zhang,† Jan O. Binder,† Pascal Hartmann,§,∥ Wolfgang G. Zeier,*,†,‡ and Jürgen Janek*,†,‡,§ †

Institute of Physical Chemistry, Justus-Liebig-University Giessen, Heinrich-Buff-Ring 17, D-35392 Giessen, Germany Center for Materials Research (LaMa), Justus-Liebig-University Giessen, Heinrich-Buff-Ring 16, D-35392 Giessen, Germany § BELLA − Batteries and Electrochemistry Laboratory, Institute of Nanotechnology, Karlsruhe Institute of Technology, Hermann-von-Helmholtz Platz 1, D-76344 Eggenstein-Leopoldshafen, Germany ∥ BASF SE, 67056 Ludwigshafen am Rhein, Germany ‡

S Supporting Information *

ABSTRACT: All-solid-state lithium ion batteries may become long-term, stable, high-performance energy storage systems for the next generation of electric vehicles and consumer electronics, depending on the compatibility of electrode materials and suitable solid electrolytes. Nickel-rich layered oxides are nowadays the benchmark cathode materials for conventional lithium ion batteries because of their high storage capacity and the resulting high energy density, and their use in solid-state systems is the next necessary step. In this study, we present the successful implementation of a Li[Ni,Co,Mn]O2 material with high nickel content (LiNi0.8Co0.1Mn0.1O2, NCM-811) in a bulk-type solid-state battery with β-Li3PS4 as a sulfidebased solid electrolyte. We investigate the interface behavior at the cathode and demonstrate the important role of the interface between the active materials and the solid electrolyte for the battery performance. A passivating cathode/electrolyte interphase layer forms upon charging and leads to an irreversible first cycle capacity loss, corresponding to a decomposition of the sulfide electrolyte. In situ electrochemical impedance spectroscopy and X-ray photoemission spectroscopy are used to monitor this formation. We demonstrate that most of the interphase formation takes place in the first cycle, when charging to potentials above 3.8 V vs Li+/Li. The resulting overvoltage of the passivating layer is a detrimental factor for capacity retention. In addition to the interfacial decomposition, the chemomechanical contraction of the active material upon delithiation causes contact loss between the solid electrolyte and active material particles, further increasing the interfacial resistance and capacity loss. These results highlight the critical role of (electro-)chemo-mechanical effects in solidstate batteries.

1. INTRODUCTION

products and low stability at elevated temperatures strongly affect the battery performance.9−11 Therefore, all-solid-state batteries (SSBs) without liquid electrolyte are studied intensively as potential competitors of LIBs. Instead of the commonly employed carbonate-based liquid electrolytes, SSBs contain a consolidated polycrystalline or glassceramic solid as the separator as well as a solid electrolyte (SE) in the electrode composites. Employing SEs is thought to solve the issues associated with liquid electrolyte LIB, due to a higher degree of safety by avoiding leakage, and concurrently high energy densities.1 In the search for feasible solid electrolytes, the lithium ion conductivity hasin a few casesreached values

Environmental concerns and the growing fraction of renewable energy in our energy supply drive the development of electric and hybrid electric vehicles with high-energy batteries. Lithium-ion batteries (LIBs) based on liquid electrolytes are expected to serve as the most reliable energy storage devices in electric vehicles. Within the next few years, these “conventional” LIBs will reach their performance limits in terms of energy density.1 Coming closer to these performance limits, safety concerns grow.2−5 Commonly employed liquid electrolytes can potentially leak or evaporate and are highly flammable as well as toxic.6 In addition, liquid electrolytes currently do not allow the use of high voltage cathodes or lithium metal anodes, due to limited stability at high potentials on the one hand and dendrite growth on the other hand.7,8 Furthermore, electrode cross-talk by diffusion of side © 2017 American Chemical Society

Received: March 7, 2017 Revised: June 8, 2017 Published: June 9, 2017 5574

DOI: 10.1021/acs.chemmater.7b00931 Chem. Mater. 2017, 29, 5574−5582

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Chemistry of Materials that are even higher than those of their liquid counterparts.1 In particular, glass-ceramic electrolytes within the system Li2S−P2S5 have high lithium ion conductivities of about 10−3 S/cm, which is enough to achieve a good solid-state battery performance.12 In addition to the high Li ion conductivity, sulfide electrolytes such as Li3PS4, Li7P3S11, or Li10GeP2S12 have favorable mechanical properties as ductility and elasticity, which even allow coldpressing.13−16 However, these thiophosphate electrolytes are not stable against lithium metal,7,8,17−19 and the chemically more stable but mechanically rigid Li7La3Zr2O12 garnet electrolyte is also considered as a potential SE candidate for SSBs.20−23 Recent efforts by Kanno and co-workers24 offer insights in the potentially high performance of SSBs using intercalation anodes. In high performance lithium ion batteries, layered oxide solid solutions like lithium nickel−cobalt−manganese oxide (LiNi1−x−yCoxMnyO2, in short NCM) or lithium nickel−cobalt− aluminum oxide (LiNi1−x−yCoxAlyO2, in short NCA) and their derivatives are mainly used as cathode materials.25,26 Compared to the original LiCoO2, the solid solutions offer higher capacities due to the high Ni content, increased safety, and potentially reduced cost.27−29 However, while NCM materials offer a capacity of up to 200 mAh/g, previous studies reported a low first cycle efficiency for the application of NCM materials in all-solidstate cells, losing roughly 30% of their initial capacity during discharge.30−32 The group of Shigematsu reported a discharge capacity of 115 mAh/g for ZrO2 coated LiNi1/3Co1/3Mn1/3O2 when using a sulfide eletrolyte.30 Tatsumisago et al. achieved a similar discharge capacity of 118 mAh/g compared to the 150 mAh/g of the initial charge.31 Although improvements in decreasing the interfacial resistance were made by amorphous buffer layers like ZrO2 or lithium titanate (Li4Ti5O12), the origin of the described capacity loss in the first cycle has not been elucidated yet.30,33 In SSBs, the solid electrolyte/electrode interfaces govern the electrochemical properties and cycle lifetime even more than in conventional lithium ion batteries, as unwanted reaction products at the interface cannot dissolve and diffuse in the SE.29,34 Especially, when using a sulfide electrolyte, the performance crucially depends on the interface of the oxide material and the solid electrolyte.35−37 In this study, we therefore present an in-depth analysis of the first cycle performance of a nickel-rich layered oxide Li(Ni1−x−yCoxMny)O2 (x = 0.1 and y = 0.1), in the following abbreviated as NCM-811, for the application in a thiophosphate-based solid-state battery. We combine the active material NCM-811 with β-Li3PS4 (space group Pmna) as the solid electrolyte and investigate the charge and discharge performance and the resulting interfacial electrochemical characteristics. The electrochemical performance is evaluated using in situ electrochemical impedance spectroscopy, as well as post mortem analysis by X-ray photoelectron spectroscopy. We show that the first cycle capacity loss is correlated to the formation of a cathode/solid electrolyte interphase (CEI), which grows significantly in the first charging cycle. In addition, scanning electron microscopy reveals severe morphological changes of the cathode microstructure suggesting contact loss between the active material and the SE. As SEs show quite limited plasticity and are unable to flow or infiltrate pores in contrast to their liquid counterpartsmechanical contact loss due to irreversible shape changes of the NCM material is expected to have a significant contribution to the observed interfacial resistance.

In contrast to previous efforts that are focused on overcoming the fundamental problems of SSBs like interfacial processes, general cell design, and mechanical properties,1,24,29,38−40 the herein presented work is aiming at the evaluation and understanding of competitive materials in high performance cells. Our results not only explain the first cycle capacity loss, they rather highlight the importance and so far underrated role of chemo-mechanical effects in solid-state batterieswhich also appear to play an important role in LIBs with liquid electrolytes.41

2. EXPERIMENTAL METHODS Cell Assembly. The crystalline solid electrolyte β-Li3PS4 and the electrode material NCM-811 were provided by BASF SE for these studies. Prior to use, the received NCM-811 powder was dried in a vacuum Büchi oven at 250 °C overnight and stored in an argon-filled glovebox (H2O < 0.1 ppm and O2 < 0.1 ppm). An indium foil (Alfa Aesar, 0.125 mm thickness, approximately capacity = 7.7 mAh/cm2) with a diameter of 6 mm was used as anode material in order to prevent decomposition and side reactions.24,29,42,43 The cell casing was manufactured in-house. Cells were assembled following the previously described procedure.29 A total of 12 mg of a composite cathode powder consisting of NCM-811 and β-Li3PS4 in a mass ratio of 70:30 (volume ratio 47:53) was used. The cathode composite was ground in an agate mortar for 15 min before assembling the cell. A mass of 60 mg of the sulfide electrolyte served as a separator, corresponding to a thickness of ∼400 μm. Powders were compressed uniaxially at 35 kN before adding the indium foil, which corresponds to an applied pressure of approximately 445 MPa. During electrochemical experiments, a reduced pressure of approximately 70 MPa was applied. A total of four cells were used, and the shown data are representative for all collected data. A similar cell was constructed using 30 mg of a composite of 35 wt % lithium titanate (Li4Ti5O12, provided by BASF SE) and 65 wt % β-Li3PS4 as a negative electrode (volume ratio 23:77). The anode composite was balanced to have a 5% capacity overhead over the positive electrode to prevent limitations of the negative electrode. Cells with liquid electrolytes for comparison were prepared as cointype cells. The working electrodes consist of 90 wt % active material, 5 wt % Super P Li as the conductive additive, and 5 wt % polyvinylidendifluoride (PVdF, Solef 5130) solved in N-methyl pyrrolidone (NMP, Sigma-Aldrich) as a binder. The electrodes were prepared as a slurry, which was stirred for 24 h to archive a homogeneous distribution of all compounds in the suspension. Afterward, the slurry was cast on fresh etched aluminum foil. The cast was dried for 24 h in the air. After drying, circular electrodes with a diameter of 12 mm were prepared. To densify the active material, the electrodes were calandered (DPM Solutions) with a pressure of ca. 0.6 MPa. The electrodes were transferred into an argon-filled glovebox, after previous drying for 24 h in a vacuum oven to remove the last traces of water and solvent. For cell preparation, CR2032 coin cells cases (MTI Corporation), one Ø 16 mm Celgard 2500 separator, and 15 μL of EC/EMC (50:50 wt/wt; SigmaAldrich) as the electrolyte were used. One molar LiPF6 (Sigma-Aldrich) as the standard electrolyte salt was employed. The cells were automatically closed with a crimper machine (MTI Corporation). Electrochemical Characterizations. Electrochemical impedance spectroscopy (EIS) was performed using the EC-Lab Electrochemistry SP300, Biologic, after 1 h of galvanostatic cycling at 214 μA/cm2 (0.1 C). Measurements were conducted at the initial OCV (approximately 0.6 V) in a frequency range of 7 MHz to 1 Hz applying a 10 mV signal amplitude. The partial resistances were fitted using RelaxIS and, to the best of our knowledge, assigned to the corresponding cell components and interfaces. Long-term charge and discharge tests were performed on a MACCOR battery cycler. Cells using indium foil as a negative electrode were galvanostatically charged to 3.70 V and discharged to 2.00 V. For cells employing lithium titanate as anode material, the voltage window was adjusted to 1.00−2.75 V. The potential of the negative anode was assumed to be 0.6 V vs Li/Li+ metal for indium foil and 1.55 V vs Li/Li+ for lithium titanate, respectively.16,44 The selected 5575


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Figure 1. (a) Representative charge and discharge profiles for the first, second, and 50th cycle (blue). The SSBs were cycled between 2.7 and 4.3 V vs Li+/Li. The orange curve shows the first charge and discharge cycle of a liquid electrolyte cell using NCM-811, showing the electrochemical behavior of the SSB in comparison to the liquid electrolyte cell. The current density for all cycles corresponds to 0.1 C. (b) Rate test and long-term cyclability for the SSB: 0.1 C, 0.25 C, 0.5 C, 1 C, followed by open end cycling at 0.1 C. Red triangles indicate the charge; blue circles represent the discharge capacity. A large first cycle capacity loss was observed in the SSB, which does not occur in the liquid electrolyte cell. test procedure includes a staircase increase of the C-rate (0.1, 0.25, 0.5, 1 C) for five cycles each, followed by at least 30 consecutive cycles at 0.1 C. Batteries were cycled in a temperature controlled environment at 25 °C. The current density was calculated based on the theoretical capacity of 200 mAh/g of the cathode material. Scanning Electron Microscopy. Microstructure images of the cathodes extracted from disassembled cells were obtained on a Merlin high-resolution scanning electron microscope (Carl Zeiss AG, Germany). X-ray Photoelectron Spectroscopy. X-ray photoelectron spectroscopy was employed to identify reaction products at the interface of the cathode and solid electrolyte of disassembled cells. Measurements were carried out using a PHI5000 Versa Probe II with an Al anode. To avoid air exposure, the pelletized samples were transferred from an argon-filled glovebox to the analysis chamber using a transfer vessel filled with argon gas. Samples were measured as obtained after disassembling, without the application of argon etching as this may affect the surface chemistry. Secondary electron imaging (SXI) was used in order to find a homogeneous spot on the sample surface. The probed surface area was 100 μm × 1400 μm (i.e., X-ray spot size), and an X-ray power of 100 W was used. The pass energy of the analyzer was set to 23.5 eV for detailed spectra and to 187.9 eV for survey scans. All spectra were charge corrected to a binding energy of 284.8 eV for the C 1s line corresponding to adventitious aliphatic carbon. Measurements were evaluated using CasaXPS V2.3.17 software. In order to minimize sulfur evaporation under ultrahigh vacuum conditions, measurements were carried out at −80 °C. A model was developed in order to fit the XPS lines and identify the newly formed species. Due to spin−orbit splitting, the S 2p and P 2p signals are each composed of a peak doublet with an intensity ratio of 1:2 and defined peak separations. The used separation was 1.21 eV for the S 2P and 0.87 eV for the P 2P lines. Pristine β-Li3PS4 is a crystalline solid electrolyte consisting of PS43− tetrahedral units. Only P−S bonds are expected, and both P−P and S−S bonds are excluded. For a reasonable fit of the S 2p line of β-Li3PS4, we needed to include two binding states. The main signal at 161.4 eV represents the P−S bond in the PS43− tetrahedra.19,45 An additional line with minor intensity may be correlated with a P−S−P type bond as in the P2S74− unit (162.7 eV). According to the model, one major Pδ+−Sδ− phosphor bond is expected, which we assigned at 131.9 eV. The second species according to the model of bridging sulfur in a P−S−P bond is expected to have a higher binding energy as the character of S is less anionic and was found at 132.7 eV, which corresponds well with the made assumptions. The utilized solid electrolyte β-Li3PS4 contains a very high ratio of the PS43−

polyhedral unit. On the basis of the total sulfur quantification, about 96 atom % is present in binding energies associated with the PS43− configuration and only a small contribution of bridging sulfur which has a slightly different binding energy. This discrepancy might be explained by an underlying amorphous phase, which was found in other crystalline lithium thiophosphate phases.46

3. RESULTS AND DISCUSSION Battery Performance. All solid-state cells (Li−In | β-Li3PS4 | NCM-811/β-Li3PS4) were galvanostatically cycled within a potential window of 2.7−4.3 V vs Li/Li+ at 25 °C. The composite cathode of the tested cells contained 70 wt % active material and 30 wt % solid electrolyte, corresponding to an areal loading of 10.7 mg/cm2. This ratio was previously found to provide optimal battery performance, ensuring sufficient lithium ion pathways as well as electronic percolation in SSB cells using LiCoO2 as the cathode material.29 Figure 1 shows representative charge−discharge curves as well as the observed capacity retention in the SSB. The first cycle charge and discharge capacities are 176 mAh/g and 124 mAh/g, respectively, resulting in a Coulombic efficiency of η = 70.5%. Each SSB cell undergoes a standard rate capability test of five cycles at 0.1 C, 0.25 C, 0.5 C, and 1 C before continuously cycling at 0.1 C up to more than 50 cycles. No conductive carbon is added to purely focus on decomposition behavior of the solid electrolyte and avoid side reactions of carbon. When increasing the charging rate to 0.25 C, the cell offers a reversible capacity of 66 mAh/g. For 0.5 C, the capacity drops to 4 mAh/g. No reversible capacity was achieved at 1 C. Higher C rates do not affect the capacity retention once switching back to lower current rates occurs. The average Coulombic efficiency at 0.1 C is 99.1%. Over the subsequent 50 cycles, the cell follows a constant capacity loss of 1−2% per charging cycle. The discharge capacity is 81 mAh/g after 50 cycles (which includes the rate capability test). It can be seen that the first cycle is crucial for the overall capacity retention, as the most severe loss in capacity occurs in the first cycle (Figure 1a). Such a strong capacity loss is not found when using NCM-811 in a liquid electrolyte cell (Figure 1a), suggesting that it is characteristic for the solid-state environment. In addition to this capacity loss, an increase in the overpotential 5576


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Figure 2. Impedance spectra recorded intermittently during galvanostatic battery cycling. Left: (a) First cycle charge and discharge profile of a Li−In | βLi3PS4 | NCM811 /β-Li3PS4 cell at 0.1 C showing current interruption corresponding to the periods of impedance measurement. Right: Impedance spectra during charge (c) and discharge periods (b). Measurements were conducted after 1 h of charging or discharging, respectively. Spectra are stacked with an offset of 40 Ω in the −Z(Im) direction. Red points indicate selected frequencies in the spectra.

(∼100 mV, see Supporting Information Figure S1a) between the first and second cycle was also observed, indicating an increase in the IR drop between the electrodes of the SSB. In order to understand the underlying reasons and evaluate the origin of this capacity loss in the first cycle, in situ electrochemical impedance spectroscopy was applied during the charge and discharge of the cell. In Situ Impedance Spectroscopy. The initial low cycle efficiency of the Li−In | β-Li3PS4 | β-Li3PS4/NCM-811 cells is further evaluated using in situ impedance spectroscopy (see Figure 2). The evolution of the different resistance contributions was monitored and correlated to the observed capacity loss. The first cycle is of particular interest as the described capacity drop is the most significant between charge and discharge. Here, the first two cycles are evaluated and fitted in order to gain a detailed understanding of the degradation processes. During the impedance measurement, the battery cell was kept at open circuit voltage. The impedance response recorded between the initial low OCV and OCV < 2.90 V could not be evaluated properly due to strong scattering in the data points at mid and low frequency and are therefore omitted here. The evolution of the impedance spectra with the OCV is shown in Figure 2 as Nyquist plots. The impedance spectra change depending on the state-ofcharge of the battery (SOC), especially in the lower frequency range. As seen in Figure 2, the applied potential differs significantly from the measured OCV, indicating formation effects at the beginning of the cycling. We assume this activation process to be the cause for the disturbance in the data, observed in measurements OCV < 2.90 V during the first charge. In order to evaluate the obtained spectra, these were divided into the high frequency (MHz range), mid frequency, and low frequency regions (10−1 Hz). An exemplary deconvolution of an impedance spectrum during the first charge at OCV = 3.00 V is shown in Figure 3. Four resistances are required for a sufficient fit of the spectrum, corresponding to an equivalent circuit of (RQ)(RQ)(RQ)(RQ) as shown in Figure 3a. The modeled RQ elements were assigned in accordance with previously suggested models by Tatsumisago et al. and by Zhang

Figure 3. (a) Bode plot of the NCM811/β-Li3PS4 | β-Li3PS4 | Li−In cell during charge (OCV = 3.00 V). Open circles indicate the measured data; solid lines represent the fit result using the shown equivalent circuit. (b) Nyquist plot of the impedance spectrum. Gray circles represent the measured data. and the black solid line indicates the fit of the data. The different resistances are illustrated as colored semicircles. RSE,bulk = volume resistance of the solid electrolyte, RSE,gb = grain boundary resistance of the SE, RSE,Cathode = interfacial resistance between the SE and the positive electrode, and RSE,Anode = interfacial resistance between the SE and the negative electrode. The colored semicircles represent a guide-to-the-eye.

et al.29,38 The high frequency semicircle unequivocally corresponds to the bulk resistance of the solid electrolyte (Rbulk, C ≈ 40 pF). The mid frequency part of the spectrum 5577


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Chemistry of Materials requires at least two RQ elements for a reasonable fit, which are assigned to the grain boundary resistance of the solid electrolyte (RSE,gb, C = 161 nF) and the cathode/SE interface resistance (RSE,cathode, C ≈ 1.1 μF). The low frequency range suggests an additional semicircle which we attribute to the negative electrode/SE interface (RSE,anode, C ≈ 0.4 mF). The fitted four resistances (illustrated in Figure 3b) from each individual spectrum are depicted as a function of the OCV, i.e., as the SOC of the cell, in Figure 4.

charging process, RSE,anode drops back into the range of the original value of roughly 40 Ω (≈ 50 Ω·cm−2). In addition to the anode interface, the interfacial resistance of the positive electrode interface undergoes a transition at a different SOC. The cathode interface resistance changes significantly during the first and second battery cycle, with a strong increase of the interfacial resistance during the first charge. This most severe effect within the first charge leads to an increase of RSE,cathode by approximately 140 Ω (≈ 180 Ω·cm−2), corresponding to an IR drop of ∼25 mV at the employed current density. We believe the difference in the overpotential estimated from the impedance spectra in comparison to the one obtained from galvanostatic cycling to be due to formation effects, as discussed above. In contrast to observations at the negative electrode interface, this observed resistance increase is almost fully irreversible and maintained in the following discharge. During the second charge, the interfacial resistance at the cathode side increases further, but here the increase is smaller than in the first cycle and is not fully translated to the subsequent charge. In order to corroborate these findings, SSB cells were constructed with Li4Ti5O12 as anode material. Similar impedance results are obtained, showing that these irreversible interfacial impedance changes are indeed effects at the cathode side (see Supporting Information) and not at the anode. The particularly strong and irreversible change of the resistance of the cathode/SE interface in the first cycle leads to the conclusion that the observed buildup of the resistance is in fact not a function of the SOC of the battery but rather a function of the applied potential. The strongest increase takes place between the OCV of 3.2 and 3.4 V. Exceeding a certain potential appears to lead to irreversible changes at the cathode interface by oxidation of the solid electrolyte. The theoretical capacity of 200 mAh/g is not fully achieved as the cutoff voltage is reached earlier due to the developed overpotential. Additionally, the solid cathode composites are not 100% dense, and thus, not every NCM particle is ionically and electronically well addressed when using two-phase composites. In the second cycle, the interface formation has already taken place and only proceeds with a much slower rate, as seen in the average charge and discharge voltages (Supporting Information). The changing RSE,cathode during the discharge process instead may originate from lithium intercalation into the NCM material. It is difficult to interpret the processes occurring during the discharge of the cell, as seen in Figure 3. This may be due to the simultaneous steep resistance increase of the low frequency semicircle, RSE,anode. X-ray Photoelectron Spectroscopy. The irreversible resistance observed by impedance spectroscopy suggests an electrochemically induced reaction of the cathode material and the solid electrolyte. In order to test this hypothesis, the NCM811 | β-Li3PS4 cathode interface was investigated using X-ray photoelectron spectroscopy (XPS). For this purpose, cathodes of disassembled cells were analyzed and compared to the cathode composite blend as well as the pristine electrolyte. The cathode composites were measured after the first charge (charged state), in which most of the observed changes occurred (Figure 5c), and after the 50th cycle (discharged state) to ensure sufficient changes induced by cycling between the selected samples (Figure 5d). Measurements of the pristine electrolyte and the NCM-811/ β-Li3PS4 blend yield an almost identical spectrum, corroborating chemical stability of the layered oxide in direct contact with the solid electrolyte. No evidence for a reaction between the pristine

Figure 4. Evolution of the four resistances obtained by fitting the impedance spectra, in the first and second cycle plotted versus the measured OCV of the Li−In | β-Li3PS4 | NCM-811 /β-Li3PS4 cell. The dotted black arrows indicate the irreversible resistance growth of the cathode/electrolyte interface in the first cycle (a) compared to the reversible increase in the second charge (b). Color code: light blue = RSE,bulk; orange = RSE,gb; dark red = RSE,Cathode; and dark blue = RSE,Anode.

The resistance evolution in Figure 4 corroborates that the SOC has different effects on the four partial resistances within the cell. The high frequency resistance is virtually independent of the SOC, which agrees well with the assumption that this resistance corresponds to transport in the bulk of the solid electrolyte (RSE,bulk). This also applies to the resistance RSE,gb assigned to the grain boundaries of the solid electrolyte. The grain boundary resistance exhibits a small increase of approximately 50 Ω within the first cycle; however, no particular dependence on the SOC can be found. We believe this effect to derive from mechanical relaxation of the solid electrolyte and the composites during cell operation because the cell components are consolidated at 35 kN (446 MPa) during battery assembly; however, for the battery cycling, a much lower constant force of 5 kN (64 MPa) is applied. The most significant changes in the impedance response appear in the processes related to the electrode interfaces, with the strongest increase found for the negative electrode interface. However, this interfacial resistance of the indium anode is only affected during the discharge of the cell. Below an OCV of 3.2 V, corresponding to approximately 80% SOC (see Figure 2a), any further discharge leads to a strong increase of the anode resistance of up to 2000% of the initial value. This behavior has recently been attributed to a kinetic hindrance using indium alloys as anodes.29 Upon further discharge, In is fully delithiated. At low Li concentrations in the alloy, the interfacial resistance increases significantly. However, during the consecutive next 5578


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species when applying an electrochemical potential on the investigated samples. Peak changes due to electrochemically induced oxidation are more emphasized in the S 2p line, rather than P 2p. Sample c was charged with 0.1 C before extraction out of the cell for the analysis and, therefore, can be correlated with the sample analyzed using impedance spectroscopy after the first charge. The majority of the S 2p signal of this sample can be explained by the species found in the pristine solid electrolyte and the composite; however, the signal has an additional shoulder at higher binding energies. In the P 2p signal, the effect is apparent, but not as pronounced as in the sulfur spectra. In the case of sulfur, binding energies in this region indicate possible formation of −S−S− polysulfide bonds. Reference XPS measurements of P2S5 (162.6 eV) and S8 (163.6 eV) underline this hypothesis17 and rule out the formation of P2S5.45 After 50 cycles of charging and discharging, the previously formed shoulders further increase in intensity, and the signals become broader. The fit model applied to the charged sample was transferred using reasonable limitations for binding energies. The amount of oxidized sulfur and phosphorus species increases between one and 50 cycles. Compared to the majority P−S binding states in the PS43− unit, the newly formed reaction products are of a more oxidized nature. Sulfide electrolytes have a very limited electrochemical stability window, which is extended due to unfavorable kinetics in the decomposition process and the overvoltage. Calculations suggest that a reduction of the solid electrolyte at the anode would only take place at lower potentials compared to the here applied potential window.7,8 At the cathode interface, the electrolyte is exposed to oxidizing conditions and forms oxidation reaction products. The resulting sulfur species appear at a binding energy of 163.7 eV, which correlates well to the formation of −S−S− bonds. This corresponds to oxidative decomposition of the β-Li3PS4 solid electrolyte in which sulfur is oxidized from formally S(−2) to a higher (more positive) oxidation state. The interface reaction can be described as a potential driven formation of a cathode/SE interphase (CEI). This additional sulfur species cannot be correlated stoichiometrically to the simultaneously formed phosphorus species, and both, therefore, are at this stage considered as oxidized species of sulfur and the oxidic phosphorus, respectively. One possibility is the formation of P−S−(S−)n like polysulfide species with varying sulfur chain lengths as seen in liquid electrolyte based Li−S batteries.47 This is supported by broadening of the signal, indicating formation of additional oxidized species close to S(0). However, binding energies of the newly formed species do no fully correlate to sulfur−oxygen environments.45 The presented model leads to minor discrepancies of around 165 eV. Considering that not only one, but rather a large variety of oxidized sulfur species may form, the signal may be better described by a continuum of signals with nondistinct binding energies. Taking the spectral resolution into account, the collected data support a simple model for the interphase formation. The measured binding energies do not agree with the formation of S−Ox; however, P−Ox species are detected, which may indicate a reaction of the NCM material with the solid electrolyte.45 The quantification of the P 2p signal yields 7.2 atom % of the species with the highest binding energy, which corresponds to approximately 1% of the solid electrolyte to have reacted. The rather small contributions of this new reactive species suggest that the bulk electrolyte is indeed still intact and

Figure 5. Peak deconvolution of the S 2p and P 2p spectra of (a) pristine β-Li3PS4, (b) the pristine mixture of NCM811 and β-Li3PS4 without any application of current or potential, (c) the cathode composite of NCM811 and β-Li3PS4 after the first charge (0.1 C), and (d) the cathode composite of NCM811 and β-Li3PS4 after 50 cycles. Blue peaks are attributed to equivalent Pδ+−Sδ− bonds. Orange peaks were correlated to bridged P−S−P in P2S74− units. The red and green peaks are assigned to more oxidized sulfur and phosphorus species. While no reaction appears to take place just upon mixing the SE and the active material, the spectra of the charged cathodes show evidence for oxidation of the thiophosphate solid electrolyte.

materials is found, although materials have been in contact for at least 48 h. In order to further demonstrate the stability of NCM811 toward the solid electrolyte, the cathode composite was investigated by impedance spectroscopy prior to cycling experiments (see Supporting Information). The blend was analyzed as prepared and after compression as in a solid-state cell and an equilibration period of at least 180 h. Although the impedance indicates growing resistance, the effect can be correlated to the mechanical relaxation of the composite. XPS spectra remain identical before and after the equilibration period (see Supporting Information). As a consequence, any additional species can be attributed to being a result of electrochemical treatment. When comparing the electrochemically treated samples (Figure 5c and d) to the mechanically mixed composite of the active material and the solid electrolyte, oxidative changes occur in both, the S 2p and P 2p line, respectively. Both signals broaden toward higher binding energies with a more pronounced shift in the sample that had been cycled 50 times at 0.1 C. The occurrence of signals at higher binding energies indicates the formation of new oxidized sulfur and oxidized phosphorus 5579


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the as-prepared sample shows a close and intimate contact between the edges of the active material and the electrolyte itself (Figure 6b). The cathode particles are well embedded in the solid electrolyte forming a dense composite. After charging, the particles of the active material are surrounded by a spherical gap (Figure 6c, zoom-in Figure 6d). After 50 charge and discharge cycles, a similar morphology was found (Figure 6e). Thus, the observed capacity fading is likely originating from the ongoing decomposition of the SE. While mechanically induced contact loss is possible after disassembling, the symmetrical shape of electrolyte shells as well as the good contact of the particles in the uncycled cells suggests otherwise. A negative imprint of the surface morphology of the NCM particles is found on all edges in the solid electrolyte surface (Figures 6d,f), indicating a former intimate contact between both materials that has been lost during cycling. The resulting contact loss upon cycling of the NCM particles with the solid electrolyte can be well explained by the volume changes of NCM materials. During delithiation, the unit cell volume of NCM shrinks,41,48 compared to LiCoO2 that undergoes an overall expansion of the unit cell.40 The volume change is especially high for nickel-rich materials as used in this study.48 We suspect that the cathode interface formation in combination with the mechanical deflation provokes a contact loss between active material and solid electrolyte at the first charge. Since the composite cathodes are compressed prior to charging, subsequent contact loss in an all-solid-state system will have severe consequences on the performance of the material, as no additional solid electrolytes can fill the emerging voids. The result is a partial electrochemical contact loss reducing the overall achieved capacity retention during the subsequent cycles. While the contact loss should only occur during the initial charge, corresponding well to the initial capacity drop, the ongoing capacity fade can be attributed to the propagating interphase formation and decomposition reaction.

that the electrochemical decomposition takes place at the interface only. Comparing the approximate amount of reaction products between the first and 50th cycle, it is evident that no steady decomposition takes place. Most of the interphase has formed after the first charging cycle, as supported by the impedance analysis (Figure 4). The increase in overvoltage for the consecutive cycles is rather small (see Supporting Information). This interphase formation, however, has severe consequences for the capacity retention of this cycle particularly and affects the cutoff voltage in the following cycles. The growing interfacial resistance RSE,cathode suggests a decreased lithium ion conductivity in the newly formed CEI, which possibly results in partial insulation of NCM particles due to the forming elemental sulfur species. Cathode Microstructure. To better understand the evolving interfacial resistance at the cathode side, the cathode composite morphology was investigated using scanning electron microscopy. In line with the XPS experiments, electrodes were scanned after the initial charge and after 50 cycles. Representative cells were disassembled and fragments of the composite cathode extracted. For comparison and elimination of any mechanical artifacts due to the disassembly of the cells, an additional battery cell of the same composition was prepared without exposing it to any electrochemical treatment. The NCM material is easily identified by its characteristic spherical morphology. Fragments of the spherical secondary particles arise from processing of the composites in the agate mortar or disassembling of the cells. Besides some minor cracks in the solid electrolyte (Figure 6a),

4. CONCLUSION We have evaluated the interfacial processes for a solid-state bulk type battery using the high capacity, nickel-rich cathode material NCM-811 and β-Li3PS4 as the solid electrolyte. The origin of the low first cycle efficiency of the solid-state battery was investigated by state-of-charge dependent impedance spectroscopy, X-ray photoemission spectroscopy, and scanning electron microscopy. From in situ impedance spectroscopy, it was observed that an additional resistance is formed irreversibly at the positive electrode interface in the first charging period, which is suspected to be responsible for the low capacity retention in the first cycle. This resistance was found to originate from a highly resistive interphase in the cathode which is formed by oxidation of the SE during the charging periods. Evaluation of the XPS data suggest that the majority of the passivating layer is developed during the first charge and grows slowly upon further cycling. Furthermore, it was found that particles of the active material lose contact with the solid electrolyte due to the chemical contraction of NCM811, which suggests that those are no longer fully electrochemically addressed during the subsequent cycles. Our results show that the observed capacity loss during the first cycle is a combination of changes in the chemical composition at the interface of the solid electrolyte (oxidation) as well as contraction of the NCM particles during delithiation (charging). This highlights the critical relevance of interfacial reactivity, cathode/ electrolyte interphase formation, as well as of (electro-)chemo-

Figure 6. Scanning electron micrographs of the cathode composite of NCM811 and β-Li3PS4 (a,b) as prepared in a solid-state cell but without the application of current or potential. (c,d) SEMs of a Li−In | β-Li3PS4 | NCM-811 /β-Li3PS4 cell after single charging to 4.3 V vs Li/Li+ at 0.1 C. (e,f) SEMs of a given cell after 50 full battery cycles in the discharged state. The NCM-811 particles shrink during delithiation (charge) and lose contact with the SE. The remaining imprint of the NCM morphology on the electrolyte shows the initially intimate contact between the active material and the electrolyte. 5580


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mechanical processes of the active materials in solid-state batteries.

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ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.chemmater.7b00931. EIS spectrum of the solid separator without electrodes, charge−discharge profiles and average charge−discharge voltages showing the development of the overpotential upon cycling, state-of-charge dependent EIS data for an SSB cell using a Li4Ti5O10 composite anode as a comparison, and additional XPS and EIS measurements to demonstrate the stability of the solid electrolyte in contact with the active material over time (PDF)

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AUTHOR INFORMATION

Corresponding Authors

*E-mail: [email protected]. *E-mail: [email protected]. ORCID

Wolfgang G. Zeier: 0000-0001-7749-5089 Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS The authors acknowledge financial support by BASF SE within the International Network for Electrochemistry and Batteries. W.G.Z furthermore gratefully acknowledges the financial support through start-up funding provided by the JustusLiebig-University Giessen. R.K. acknowledges support by the Funds of the Chemical Industry (FCI). The authors thank Dr. Bjoern Luerßen for his contribution to the abstract figure.

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REFERENCES

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Capacity Fade in Solid-State Batteries: Interphase Formation and

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