Understanding the Capacity Loss in LNMO-LTO Lithium-Ion Cells at

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C: Energy Conversion and Storage; Energy and Charge Transport

Understanding the Capacity Loss in LNMO-LTO LithiumIon Cells at Ambient and Elevated Temperatures Burak Aktekin, Matthew James Lacey, Tim Nordh, Reza Younesi, Carl Tengstedt, Wolfgang Zipprich, Daniel Brandell, and Kristina Edstrom J. Phys. Chem. C, Just Accepted Manuscript • DOI: 10.1021/acs.jpcc.8b02204 • Publication Date (Web): 03 May 2018 Downloaded from http://pubs.acs.org on May 5, 2018

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The Journal of Physical Chemistry

Understanding the Capacity Loss in LNMO-LTO Lithium-ion Cells at Ambient and Elevated Temperatures Burak Aktekin1, Matthew J. Lacey1*, Tim Nordh1, Reza Younesi1, Carl Tengstedt2, Wolfgang Zipprich3, Daniel Brandell1, Kristina Edström1 1

Department of Chemistry – Ångström Laboratory, Uppsala University, Box 538, SE-75121, Uppsala, Sweden 2 Scania CV AB, SE-151 87, Södertälje, Sweden 3 Volkswagen AG, D-38436, Wolfsburg, Germany *[email protected]

Abstract The high voltage spinel LiNi0.5Mn1.5O4 (LNMO) is an attractive positive electrode due to its operating voltage around 4.7 V (vs. Li/Li+) and high power capability. However, problems including electrolyte decomposition at high voltage and transition metal dissolution, especially at elevated temperatures, have limited its potential use in practical full cells. In this paper, a fundamental study for LiNi0.5Mn1.5O4 || Li4Ti5O12 (LTO) full cells has been performed to understand the effect of different capacity fading mechanisms contributing to overall cell failure. Electrochemical characterization of cells in different configurations (regular full cells, back-to-back pseudo-full cells and 3-electrode full cells) combined with an intermittent current interruption technique have been performed. Capacity fade in the full cell configuration was mainly due to progressively limited lithiation of electrodes caused by a more severe degree of parasitic reactions at the LTO electrode, while the contributions from active mass loss from LNMO or increases in internal cell resistance were minor. Comparison of cell formats constructed with and without the possibility of cross-talk indicate that the parasitic reactions on LTO occur because of the transfer of reaction products from the LNMO side. The efficiency of LTO is more sensitive to temperature causing a dramatic increase in the fading rate at 55 ˚C. These observations show how important the electrode interactions (cross-talk) can be for the overall cell behaviour. Additionally, internal resistance measurements showed that the positive electrode was mainly responsible for the increase of resistance over cycling, especially at 55 ˚C. Surface characterization showed that LNMO surface layers were relatively thin when compared to the SEI on LTO. The SEI on LTO does not contribute significantly to overall internal resistance even though these films are relatively thick. XANES measurements showed that the Mn and Ni observed on the anode were not in metallic state; the presence of elemental metals in the SEI is therefore not implicated in the observed fading mechanism through a simple reduction process of migrated metal cations.

Introduction LiNi0.5Mn1.5O4 (LNMO) is a promising spinel-type positive electrode (cathode) material for lithiumion batteries (LiBs), with a theoretical capacity of 147 mAh/g and an operating voltage around 4.7 V (vs. Li/Li+). This high voltage plateau, due to the active Ni2+/Ni4+ redox couple, renders a high theoretical energy density of 690 Wh/kg for the active material.1 Additionally, the intrinsic properties of the spinel structure allow excellent rate capability, and combined with the comparatively high energy density, these two main advantages make the material an ideal candidate for high power applications (e.g., electric vehicles).2 However, the anodic instability of conventional LiB electrolytes (namely LiPF6 in organic carbonates) at high voltages cause interfacial side reactions which are also accompanied by transition metal dissolution. This, in turn, causes rapid capacity fading in full cells, especially at elevated temperatures.3,4 ACS Paragon Plus Environment

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It is important to understand how undesired reactions originating from the cathode side affect the aging of different cell components and how these contribute to the overall cell life. If classified into three main groups, the different contributions to capacity fading can be due to ‘active mass loss’, ‘internal resistance increase’ of the cell and ‘loss of cyclable lithium’.4 Active mass loss and internal resistance increase can result from the degradation of the active material or other electrode components such as the binder, conductive additive and current collector. Similarly, oxidation of the electrolyte at high potentials and interactions between the positive and negative electrodes might influence the cell life-time significantly.5 In optimized LNMO-based composite electrodes, the electrode integrity remains relatively intact, as the volume expansion of LNMO during cycling is moderate (around 7%)6 and binders such as PVdF7 have acceptable anodic stability. Conductive additives such as carbon black have been shown to be reactive8,9 and the Al current collector has been shown to dissolve to a minor extent10, but only significantly at high potentials (>5 V vs Li/Li+) or after long-term cycling: for these reasons, the degradation of inactive electrode components in the LNMO electrode is not expected to have a major role in active mass loss or internal resistance on the shorter timescale of most published studies in the literature. On the other hand, transition metal dissolution, which is more severe at elevated temperatures,11 can cause active mass loss from the cathode. However, the contribution of metal dissolution to capacity fading directly as a result of active cathode mass loss has been reported to be less than 5%, as confirmed by the recovery up to 95% of the capacity in a LNMO half cell constructed with an aged electrode from a LNMO-graphite full cell.12 This verifies that ‘the active mass loss’ due to metal dissolution (or electrode degradation) has only a minor contribution to the observed rapid fading of LNMO based full cells during these limited testing times. Capacity fading caused by internal resistance increase can also be described as power fading, since it affects the cell capacity due to kinetic limitations for a specific test or usage condition. For instance, the generation and precipitation of organic compounds, LiF, C-F and P-Fx species on LNMO might result in sluggish Li+ intercalation/deintercalation kinetics13 and therefore might increase the electrode impedance.14 Apart from this, interactions between the electrodes (“cross-talk”) are known to occur in LNMO based full cells. Migration of oxidation products from the positive electrode can result in the formation of surface films on the negative electrode,15 and migration of dissolved metal cations can cause damage to an existing Solid Electrolyte Interphase (SEI) on the negative electrode, resulting in further thickening of the SEI.11 Further contributions may include the evolution of gas products such as CO2 after electrolyte oxidation16 that can lead to increased cell resistance17; or clogging of electrode and separator pores by oxidation products.4,18 In the study reported by Kim et al.,12 electrodes were harvested from cycled LNMO-graphite full cells which had shown significant capacity loss and reassembled into half cells: both half cells were largely able to recover their initial capacities, indicating that the capacity loss in the full cell was dominated by a third mechanism, which can be expressed as ‘loss of cyclable lithium’. Loss of cyclable lithium is an issue related to capacity balance between anode and cathode, and it is obviously not a problem in half cells in which the lithium electrode provides an effectively unlimited capacity. Such a capacity fading mechanism can occur via electrolyte reduction at the negative electrode during SEI formation19 or oxidation at the positive electrode. The result of a greater extent of one of these processes over the other is excess charge stored by one of the electrodes, which perturbs capacity balance.19 Moreover, one explanation for the capacity loss in LNMO-graphite full cells has been made by an excess of electrolyte reduction as a result of SEI damage caused by migrating metal ions from the positive electrode.11 Interestingly, when LNMO was coupled to Li4Ti5O12 (LTO), which is known for its good electrochemical stability in standard lithium-ion battery electrolytes (electrode potential around 1.5 V vs. Li/Li+) and thereby limited SEI layer formation and high coulombic efficiency, the capacity fading was found to occur by the same mechanism.15,20 This shows that crosstalk between the electrodes – not necessarily limited to dissolved metal ions – exists and has a significant role in determining cell lifetime.

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The well-known example of dissolved metal ion (e.g. Mn) interactions between electrodes has been studied extensively21 but the exact mechanisms leading to the cyclable lithium loss is not yet clear. Similar ambiguities also exist regarding the oxidation state of transition metals deposited on the negative electrode.22 Recently, some studies have reported a different kind of cross-talk: the interaction of gas products evolved in systems such as LNMO-graphite23 and NMC-graphite.24,25,26 The main gas product of electrolyte oxidation was identified as CO2 at high cathode voltages16 while gases such as C2H4, CO and H2 might evolve on the negative electrode at potentials below 1 V (vs. Li/Li+).24 Xiong et al. showed CO2 evolution when de-lithiated (charged) NMC was simply stored in the electrolyte at 60 ˚C, while no considerable gas was present after storage in charged full cells (with a graphite electrode), indicating CO2 consumption by the negative electrode.26 The reduction of CO2 to species such as lithium oxalate or lithium carbonate would cause cyclable lithium loss – and this loss would be reversible if soluble oxalate can migrate back to the positive electrode and be re-oxidized (i.e., a redox shuttle mechanism).27 In addition to gas-consuming reactions, gas-generating reactions occurring via cross-talk are also possible. Metzger et al. for example observed persistent H2 evolution at a graphite electrode caused by the reduction of protic species formed continuously by electrolyte oxidation at a NMC cathode.24 The correlation between different types of electrode interactions and the overall fading of cell capacity indicates that these are not independent processes, and more research is necessary to identify their exact roles in full cells containing high voltage electrode materials. In this paper, we aim to understand how different components of the cell contribute to capacity fading, according to the three categories of fading mechanism outlined above, in LNMO-LTO full cells cycled both at room temperature and elevated temperature (55 ˚C). As seen in the literature, the fading rate for this type of cell at elevated temperature is much faster (e.g. at 55 ˚C). In some reports, this performance difference is more pronounced; e.g., quite stable cycling for >1000 cycles has been reported at room temperature28 while fast fading was observed at 55 ˚C. We investigate the reasons behind these observations by following independent internal resistances and electrode potentials with respect to LNMO and LTO in three-electrode cells which have been cycled galvanostatically with short-duration intermittent current interruptions.29 Pseudo-full cells (i.e., ‘back-to-back’ cells) as suggested by Li et al.20 have also been tested to observe cycling behaviour when no cross-talk is possible. Regular full cells and half cells have been tested for reference as well as for the ex situ characterization of cycled electrodes to account for the effect of metal dissolution and surface layer growth. We believe that the findings from this combination of different cell configurations are highly valuable for providing insights into the ultimate failure mechanisms of ‘high voltage positive electrodes’, since LTO has a high electrode potential and does not contribute to extensive SEI formation (as opposed to for example graphite, which consumes cyclable lithium during initial film formation and film repair mechanisms following SEI damage from dissolved metal ions).

Experimental Preparation of electrodes & pouch cells. LNMO slurry was prepared via 2 hours ball milling of commercial LNMO powders (90 wt%) with carbon black (Imerys, C65, 5 wt%) and a poly(vinylidene difluoride)-based binder (5 wt%, PVdF-HFP, Kynar Flex 2801) in N-methyl-2-pyrrolidone (NMP, VWR Chemicals). The binder was pre-dissolved in NMP in 4:96 wt% ratio. The slurry was casted onto Al foil (20 µm thickness) and subsequently calendered to a final coating film thickness of approximately 50 µm. The average porosity was 40-50% and the active mass loading was around 10.6 mg/cm2 (or 1.56 mAh/cm2). Commercial LTO electrodes were provided by Leclanché and had a capacity of approximately 1.7 mAh/cm2. Therefore, the capacity of full cells were always limited by the LNMO electrode. All electrodes were dried for 10 hours under vacuum at 100 ˚C. A schematic of a standard pouch test cell is shown in Figure 1a. 2 cm diameter electrodes were punched from LNMO and LTO sheets and 2 layers of Celgard 2500 separators with 4 × 4 cm2 area were used. This separator is a monolayer polypropylene with 25 µm thickness and 64 nm average pore size. The separators were dried overnight under vacuum at 70 ˚C prior to use. 1 M LiPF6 (Ferro Corp.) dissolved in EC and DEC (BASF) with 1:1 volume ratio was used as electrolyte. The electrolyte volume was fixed at 120 µl. The ACS Paragon Plus Environment

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connecting tabs were shaped with 2 cm diameter circular ends inside the cell to ensure homogeneous pressure distribution and good matching of electrode position. Al foil was used to contact LNMO and Cu for LTO electrodes. Nickel strips were used to contact lithium metal electrodes. Pouch cells were assembled in an Ar-filled glove box with H2O < 5 ppm and O2 < 1 ppm.

Figure 1. Schematic view of pouch cell designs used for cell testing in this study. In three-electrode cells (Figure 1b), concentric lithium rings were punched with an inner diameter of 2.2 cm and outer diameter of 2.6 cm and carefully placed between LNMO and LTO, facing separators on both sides. In half cells, the lithium counter electrode diameter was 2.6 cm. Pouch cell size was 10 × 5 cm and during cycling, half of the pouch (5 × 5 cm) was compacted using flat plates to ensure better electrical contact. The extra 5 × 5 cm area was for accumulation of evolving gases in order to reduce the risk of cell leakage or bubble formation in pores. So-called “back-to-back cells” (Figure 1c) were constructed by connecting the working electrodes of separate LTO and LNMO half cells as the negative and positive terminal respectively, and shorting their respective lithium electrode terminals. The lithium electrode terminals were connected to the reference electrode terminal on the potentiostat to allow for separate measurements of the LTO and LNMO electrodes. A modified back to back cell was also constructed by using only one pouch cell. This time, 4 sheets of separator were used and a large piece of nickel metal foil covered with 2.6 cm diameter Li discs on both sides was placed between the LNMO and LTO electrodes so that the lithium metal faced two separators on both sides. There was no direct electrolyte contact between the two sides of this foil (Li+Ni+Li), however, evolving gases from one electrode would not be prevented from coming into contact with the other.

Electrochemical characterization. All cells were galvanostatically cycled between cell voltages of 1.5 V and 3.5 V (i.e., vs. LTO) at C/5 rate according to the calculated theoretical capacity of LNMO. In all ACS Paragon Plus Environment

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measurements, the cells were kept under OCV condition for 10 hours at the testing temperature prior to cycling. A Novonix HPC (High Precision Charger) system was used for testing of regular (two electrode) cells at 55 ˚C while a Digatron BTS-600 was used for room temperature (24 ˚C – 25 ˚C) testing. Testing of three electrode and back-to-back cells was made via galvanostatic cycling combined with an intermittent current interruption (ICI) resistance determination method29,30 using a Bio-Logic VMP2 instrument. Cells were tested at room temperature, 40 ˚C and 55 ˚C. The current was interrupted after every 150 seconds for 1 second. The calculation of a time-independent cell internal resistance was made by regression of E vs. t1/2 to t=0 for the interruption period, and calculated according to Ohm’s law (i.e., R = ∆E/∆I). Data analysis was performed using the R programming language and environment. Simple simulations of capacity fade due to coulombic efficiency imbalance were also written in R; the code is supplied in the Supporting Information (SI).

Surface characterization. Analysis of pristine LNMO and LTO electrodes were performed prior to any electrolyte contact, and on electrodes after cycling at C/5 rate for 50 cycles both at room temperature and at 55 ˚C. Electrode morphologies were imaged via a Zeiss 1550 scanning electron microscope (SEM). Cycled pouch cells were opened in an Ar-filled glove box. LNMO and LTO electrodes were rinsed with DMC five times using 4-5 droplets each time to remove salt residues. Electrodes were placed on carbon tapes after drying off the DMC. Samples were transferred to the SEM in airtight glass vials and exposed to air for 15-20 seconds before being transferred to the SEM chamber. The accelerating voltage was 3 kV and the working distance was 5 mm during analysis. All images were scaled to the same magnification using the image analysis software ImageJ. Surface characterization was made via X-ray photoelectron spectroscopy (XPS). The sample preparation was the same as for SEM except an airtight transfer system was used for sample to avoid any exposure to air. All samples were analyzed using a Phi-5500 instrument with monochromatized Al-Kα radiation (1486.6 eV). Additionally, LTO electrodes were analyzed via HAXPES for depthprofiling of surface films. To ensure reliability of results from the different analyses, multiple samples were cut from each electrode and vacuum sealed into coffee bags (two times) using protective boxes. Measurements were performed at the Helmholtz Zentrum Berlin, BESSY KMC-1 beamline, High Kinetic Energy Photoelectron Spectrometer (HIKE) end station using a Scienta R4000 electron analyzer. Sealed samples were opened in an Ar-filled glove box, mounted on a sample holder with copper adhesive tape and transferred to the analysis chamber without any air exposure. Photon excitation energies were 2005 and 6015 eV. XANES measurements were performed at the same beamline using a Bruker XFlash 4010 fluorescence detector. Data calibration was made by linear shifting of the hydrocarbon peak to 285 eV for the LNMO data. The XPS data for the LTO electrodes were calibrated using the titanium (Ti 2p) peak. The titanium peak was set to be 459 eV as measured previously for LTO with carboxymethyl cellulose (CMC) binder.31 For the 1486.6 eV measurement of the sample cycled at 55 ˚C, no Ti 2p peak was visible, therefore the main peak in the Mn 2p spectra was adjusted to align with the corresponding peak observed with HAXPES data (2005 eV) as shown in supplementary information (SI). This shift was applied to all other spectra in that sample. All normalization was done by normalizing the area of the spectra to 1. For measurements with no peaks, the noise level was adjusted to match instrument noise level for other measurements. For XANES, energy calibration was made using gold foil mounted on the sample holder. The monochromator was adjusted to the edge energy and Au 4f spectra were measured to determine the energy shift. CasaXPS and Igor were used for the analysis of XPS and HAXPES data. Analysis and normalization of XANES data were made using Athena software.

The effect of positive and negative electrode efficiency on capacity fading: a simple model As mentioned in the introduction section, half cells with optimized LNMO electrodes and a Li counter-reference electrode with a non-limiting capacity are able to operate without significant ACS Paragon Plus Environment

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capacity loss, while capacity loss is apparent in full cells where both electrodes have a limiting lithium inventory: this is a clear indication of cyclable lithium loss during cycling. It is necessary at this point to clarify certain terms. The term ‘lithium inventory’ refers to the amount of lithium hosted in an electrode at any given time; in the absence of side reactions, the sum of the lithium inventory of both electrodes is constant. On the other hand, the term ‘cyclable lithium’ refers to the usable (or electrochemically accessible) lithium content of the cell, and is distinct from the sum of the lithium inventory of the constituent electrodes. A loss in cyclable lithium is observed both at room temperature and elevated temperatures20 in LNMO-based full cells. It is important to understand the reasons that lead to such a loss to meaningfully interpret the response of cells to, for example, different cycling conditions, and electrolytes.32 We can consider three extreme cases with simple simulations of arbitrary full cells in which different ‘coulombic efficiencies’ are assigned to each electrode. We can assume a larger degree of parasitic reactions during charging compared to discharging, on both electrodes, due to overpotential: for the purpose of this exercise we assume that parasitic reactions take place only during charge, and that the discharge is 100% efficient for both electrodes. Simulations were made for LNMO-limited full cells (i.e., excess capacity in LTO) for comparison with the cells tested experimentally in this work (an example of the code is given in the Supporting Information). In the simulations, we assigned initial capacities such that the ratio of the electrode capacities are comparable with the experiments presented in this work. A capacity of 5 (“mAh”) was chosen for the negative (LTO) electrode and 4 (“mAh”) for the positive (LNMO). For simplicity, it is assumed that the electrode impedance does not affect the cycling performance and no active mass loss occurs. First, we observe how the cyclable lithium loss would be manifested in full cells if parasitic reactions are limited to the LNMO electrode (see Figure 2a-b). Consider a case where LNMO has an efficiency of 0.95 (i.e., 95%) and LTO an efficiency of 1. That is, for every 1 mole of charge passed through the cell, the lithium inventory in LNMO changes by 0.95 moles, and 0.05 moles of charge is passed through a parasitic reaction at the positive electrode. In Fig. 3b, the black line shows the instantaneous amount of charge stored (Li+ ions, i.e., lithium inventory) in LNMO (blue for LTO). LTO has zero initial lithium inventory, LNMO is, correspondingly, initially fully lithiated. As clearly seen in Figure 2b, a parasitic process at the LNMO electrode effectively extends the charging time and results in the intercalation of a greater amount of lithium into LTO (~4.2 mAh) than is initially stored in LNMO (4 mAh); the excess Li+ is drawn from the electrolyte. However, this “excess” lithium in LTO cannot later intercalate back into LNMO during discharging since the capacity of LNMO is still limited to 4 mAh. Therefore, nearly 0.2 mAh of lithium inventory remains ‘trapped’ in LTO after one complete cycle. It is seen in Figure 2a that, in this scenario, the overall cell capacity is stable over the first 5 cycles but thereafter fades at an almost constant rate. The stable early cycling results from the excess capacity in anode – allowing complete delithiation of the positive electrode, but for only a limited time. It is clear from Figure 2b that the “active cycling window” shifts to the higher lithiation states of both electrodes during cycling. From the 5th cycle onwards, the lithium inventory in the negative electrode reaches the full 5 mAh capacity at the end of charging step. Complete delithiation of LNMO is therefore no longer possible during the subsequent cycles. Capacity fading is observed from this point because the parasitic process becomes progressively limiting, and the cycling window of both electrodes becomes progressively more limited to the higher lithiation states. Therefore, ultimately, LNMO and LTO would both be ‘trapped’ in their fully lithiated states and the amount of cyclable lithium will become zero. In this type of failure mechanism, the cyclable lithium loss need not be attributed to any chemical inactivation of lithium in the system (as compared to, for example, formation of inert lithium-containing compounds) but rather an accumulation of excess lithium in the negative electrode. In this scenario, the coulombic efficiency of the cell is equal to the coulombic efficiency of the positive electrode. We can describe this fading mechanism as “limited delithiation”, as the lithium inventory which can be extracted from either electrode during the respective delithiation process progressively decreases. ACS Paragon Plus Environment

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Figure 2. Simulation of capacity fading with different ‘efficiencies’ of negative and positive electrodes (a, c, d). Corresponding charge amounts in both electrodes are shown below the each graph (b, d, f). In these graphs, 4 mAh refers to fully lithiated LNMO and 5 mAh to fully lithiated LTO.

In Figure 2(c, d), a different scenario is presented where a parasitic process occurs only on the negative electrode (corresponding to an LTO efficiency of 0.95; an unrealistic scenario for LTO electrodes in a standard electrolyte, but still a useful comparison). In this case, the positive electrode can be completely delithiated during charging, but since some of the charge passed is consumed in side reactions on the negative side, a smaller amount of lithium is intercalated into the negative electrode (cf. self-discharge of the negative electrode). Ultimately, both electrodes would be trapped in their fully delithiated states where the amount of cyclable lithium again approaches zero. Here, the overall coulombic efficiency of the cell is determined by the coulombic efficiency of the negative electrode. We can correspondingly describe this fading mechanism as “limited lithiation”, as the lithium inventory which can be inserted into each electrode during the respective lithiation process progressively decreases. In Figure 2(e, f), we consider a scenario where parasitic processes occur on both electrodes to the same degree (i.e., they both have 0.95 efficiency). Here, the excess charge passed by the positive electrode during charging is now balanced by an equivalent amount of excess charge passed at the negative electrode. If the efficiencies of both electrodes are equal, no imbalance in lithium inventory is created between the electrodes, and no capacity fade occurs as a result. Even though stable cycling is observed, it should be noted that this scenario still represents the worst case for the real system. Assuming the side reactions are irreversible, then electrolyte degradation (and the corresponding impedance increase) would eventually degrade the performance of the cell. Although the above scenarios represent extreme cases, parasitic reactions do occur on both electrodes in practice. We can therefore expect behavior corresponding to one of the cases depicted here, determined by the more inefficient electrode. This is true not just for the LNMO-LTO system ACS Paragon Plus Environment

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considered here but for any comparable battery system, both in operation as well as during storage (an illustrative example of these mechanisms during self-discharge is given in the Supporting Information, Figure S1). It has been observed in the literature that the cyclable lithium loss in LNMO-based full cells is due to the “limited lithiation” mechanism previously discussed (i.e., the scenario depicted in Figure 2c-d). For example, Kim et al. 12 demonstrated this mechanism in LNMO-graphite cells by further cycling of half cells made with LNMO and graphite electrodes harvested from exhausted full cells. Following the scenarios presented above, the negative electrode must therefore be less efficient than the positive electrode. Where a passivated negative electrode such as graphite is used, this inefficiency implies that the passivation layer (SEI) should therefore somehow be damaged by species generated at the positive electrode, or that the SEI is permeable to these species. Moreover, Li et al. 20 studied LNMO-LTO pseudo- (back-to-back) full cells alongside conventional full cells to show that cross-talk exists, as LTO in isolation cannot be responsible for the high negative electrode inefficiency: products of electrolyte oxidation at the positive electrode can be reactive towards the negative electrode. These scenarios described above serve as a foundation for the interpretation of our experiments here, where different cycling conditions also have been investigated.

Results and Discussion Electrochemical Cycling Performance of 2-electrode full cells. LNMO-LTO full cells cycled at C/5 rate and at room temperature showed quite stable cycling while the fading rate was much faster when similar cells were cycled at 55 ˚C at the same C-rate (Figure 3). After 50 cycles, another cycle at C/20 rate was also tested for both cells and showed only a small improvement (< 3%) in cell capacity, indicating that the change in cell resistance had only a limited effect on the observed capacity. For the 55 ˚C cell, the changes in voltage curves between the first and 50th cycle were especially apparent for regions of high LNMO lithiation, verifying cyclable lithium loss through the “limited lithiation” mechanism. The coulombic efficiency for room temperature cycling reached around 0.997 at the 20th cycle while it was around 0.98 for the 55 ˚C cell. When the two last cycles are considered (RT), the coulombic efficiency decreased from 0.998 to 0.993 when the C-rate changed from C/5 to C/20. The corresponding decrease at 55 ˚C was from 0.98 to 0.93. Both observations show that the degree of side reactions is time dependent, rather than dependent on cycle number.


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Figure 3. (a) Galvanostatic cycling of LNMO-LTO full cells at room temperature and 55 ˚C. Smaller symbols refer to coulombic efficiency, and larger symbols refer to charge (hollow) and discharge (filled) capacity. (b) Selected voltage curves from these cells. Although the difference in cycling performance for the two temperatures is large, the difference can be even more dramatic if evolving gases disrupt the electrolyte-electrode surface (e.g. pore blocking, etc.) or cause leakage of the cell. Such a rapid failure would limit the cycling duration over which a meaningful interpretation of results and ex-situ characterization of electrodes can be made. We have observed this to be a significant problem at elevated temperature: we therefore intentionally add extra space in the pouch cell construction to delay any effect of gas evolution (see experimental section). The performance of a cell prepared in an identical manner to the standard cells used in this work, but without extra buffer space in the pouch cell, is shown in Figure 4. The rapid cell failure seen in Figure 4 shows that the fast capacity fading at elevated temperature reported by a number of studies might be partially due to experimental conditions such as cell design. It also shows the importance of cell inspection (e.g. for leakage) after certain durations during cell testing.

The effect of waiting time (at OCV) in the charged state is also shown in Figure 4. The waiting time is applied at the end of the charging step (i.e. after reaching 3.5 V) and it is gradually increased from zero to 48 hours to observe the effect of parasitic reactions under static high voltage conditions. As the waiting time increases, the discharge capacity is expected to decrease significantly due to selfdischarge. If the difference between charge and discharge capacities are compared for two subsequent cycles (with different OCV time, e.g. 25th and 26th cycle), it is seen that self-discharge rate varies between 0.25-0.35 mAh g-1 h-1. A sudden change in charge capacity can be observed when the waiting time is 48 hours. Over this duration, the rate of self-discharge is increased and overpotential becomes significant (see also Figure S2 in the Supporting Information). This indicates that the cell reaches a break-down point where other contributions apart from cyclable lithium loss start to have a significant ACS Paragon Plus Environment

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role. One such example is likely to be increased internal cell resistance. The observations in Figure 4 demonstrates that rapid cell failure can eventually happen in any cell with the same chemistry, although the cell lifetime until this failure is highly dependent on experimental conditions including cell design. Therefore, it can change significantly between different studies, which can explain some of the controversies in literature. Our focus here remains in the region before the onset of such a rapid failure, and the contribution of different fading mechanisms are therefore investigated using back-toback cells and 3-electrode cells in the following sections.

Figure 4. Galvanostatic cycling results of LNMO-LTO full cells at 55 ˚C. Red symbols refer to ordinary cells used in this study, but paused in charged state with varying durations (as showed on top of the graph). Black symbols refer to another cell prepared similarly but without leaving extra space in the cell for evolved gases (without OCV in charged state). Large symbols are capacity while small symbols are Coulombic efficiency.

Galvanostatic Cycling in Back-to-back Cells. This type of full cell has been constructed by connecting one LNMO half cell to another LTO half cell by their negative poles (lithium ends) as previously proposed in reference.20 In such a pseudo-full cell, no chemical interactions are possible between the electrodes and it is chosen to see the mechanisms leading to cyclable lithium loss as well as the rate of fading when any cross-talk between electrodes can be excluded. As a modification of the test conducted by the reference paper,20 the lithium terminals of the half cells were here used as a reference electrode. Cycling of these cells was conducted using an intermittent current interruption (ICI) technique to follow individual cell resistances in both half cells throughout the cycling. The cycling behaviour at room temperature, 40 and 55 ˚C and comparison with our previously described simulations are shown in Figure 5.


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Figure 5. The capacity fading in back-to-back cells cycled at different temperatures. Inset: expected evolution of capacity according to simulation, where LNMO efficiency is indicated and LTO is assumed to be 100% efficient.

It is seen in Figure 5 that the capacity fading rate is quite fast compared to ordinary full cells. In the inset, simulation of capacity fading is also shown. In this simulation, LTO efficiency is assumed to be 1 while LNMO efficiency is either assumed to be 0.997 or 0.98 (values obtained experimentally in Figure 3). As the inset shows, the experimental fading rates in back-to-back cells are comparable to the simulation. The second important observation from these cells is the initial stable period for room temperature cycling. A similar short period is also seen at 40 ˚C, while it is entirely missing at 55 ˚C. As in the scenario depicted in Figure 2(a-b), the extra capacity of LTO (10-15% excess) is able to compensate the electrode imbalance caused by LNMO self-discharge during cycling – but only for a limited time. It is seen that this time is only limited to the formation cycles at 55 ˚C. It should be noted that a constant coulombic efficiency is assumed in the simulations, while the real cells showed lower coulombic efficiency during the initial cycles (Figure 3a). The change in individual electrode voltages and internal resistances are shown in Figure 6. The change of median resistance with cycling is also shown in Figure S3. It is clearly seen from the LNMO voltage profiles that the active cycling window of LNMO shifts from the high voltage end, meaning that the cycling is becoming more and more limited to the higher lithiation state (i.e. lower voltages). Individual LTO voltages show that the voltage drop at the beginning of charging disappears, while a charge end-point appears later, showing that the active cycling window of LTO also shifts to higher lithiation states. This constitutes the “limited delithiation” mechanism discussed previously. The shift of LNMO voltages explains the slightly better cycling of the 55 ˚C sample as compared to the simulation (Figure 5), since the cycling of LNMO is limited to the low voltage parts which reduces the degree of side reactions. After cycling and with the back-to-back cell in a fully discharged state, the constituent half cells were separated and cycled further under the same conditions. The results are given in the Supporting Information, Figure S5. Both electrodes showed reversible capacities close to their initial capacities; furthermore the LTO electrode was found to be partially charged, with the LNMO electrode in its fully discharged (lithiated) state. This further confirms the limited delithiation mechanism. The measured resistances for each electrode in this cell construction include considerable contributions from the lithium electrodes, so caution must be exercised in their interpretation. The contribution of the lithium electrode is primarily manifested as relatively high resistance in the initial cycles and an increase in resistance during the lithium stripping process; the latter is most clearly seen in the LTO compartment in Figure 6. Both electrode resistances show comparable and minor increases ACS Paragon Plus Environment

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during cycling at room temperature. Based on this limited increase, as well as the observations of individual voltages showing no considerable increase in overpotential, it can be assumed that resistance increase does not have any major effect on the capacity fading observed. At elevated temperature, lower resistance is observed for both electrodes in the formation cycle due to improved kinetics at higher temperature. However, a gradual increase of resistance for LNMO is observed while the resistance at the LTO electrode remains almost constant (see Figure S3). It can therefore be concluded that the resistance increase mainly comes from the LNMO part and it is more pronounced at high temperature, but this resistance increase seems not to be large enough to dominate the cycling behaviour in the back-to-back cells.

Figure 6. Individual voltage curves together with resistance measurements throughout the cycles of LNMO-LTO back-to-back cells at room temperature (RT), 40 and 55 ˚C for (a) the first cycle, (b) the 10th cycle, (c) the 30th cycle and (d) the 50th cycle. Galvanostatic Cycling in Three-electrode Cells. In these cells, circular lithium rings were placed as reference electrodes concentric to the circular LNMO and LTO electrodes. As seen in Figure 7, when electrode interactions are allowed, the capacity retention is much more stable compared to the back-toback cells, especially at room temperature. As demonstrated in the back-to-back cells, there is a significantly higher degree of side reactions on LNMO compared to LTO. In a regular cell, these side reactions on LNMO should still be present. However, if fading is not observed, this implies that ACS Paragon Plus Environment

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migration of reaction products to LTO causes additional side reactions which are not observed when LTO is only in contact with the standard electrolyte. If the degree of additional side reactions on LTO is comparable to LNMO, this would give an overall result similar to the third scenario described in Figure 2e-f. Additionally, when the relative increase of fading rate with increasing temperature is compared between the cell formats, it is much higher in regular full cells. This shows that the temperature dependence of the individual coulombic efficiencies are different for LNMO and LTO electrodes, the cycling stability at 55 ˚C would otherwise be comparable to RT cycling. This can explain why the fading behaviour does not follow an Arrhenius type trend in regular full cells.

Figure 7. Capacity fading in three-electrode LNMO-LTO full cells at room temperature, at 40 and 55 ˚C.

More insight into the fading mechanisms can be gained from individual electrode potentials and internal resistance measurements (see Figure 8). As seen from the voltage curves, the cycling window is progressively shifting to the lower LNMO lithiation states as the cycling proceeds. This represents the “limited lithiation” mechanism discussed previously, and in contrast to the mechanism observed in the back-to-back cells. This implies that the LTO electrode in this case is more inefficient (i.e., has a lower coulombic efficiency) than the LNMO electrode. At room temperature, the capacity loss is relatively small, and the shift of LNMO voltage is not easily seen; however, it is clearly seen that the voltage drop at the beginning of charging is persistent for LTO. This shows that LTO is always returning to its initial delithiated state at the end of discharge and that the coulombic efficiency should be dominated by the higher degree of inefficiency on the LTO side (similar to scenario in Figure 2d). The shift of the LNMO cycling window is very clear at 55 ˚C. Similar to the back-to-back cells, the three-electrode cell was disassembled in the discharged state with the LNMO and LTO electrodes reassembled into new half cells (Supporting Information, Figure S5). In this case, the LTO electrode was found to be in its fully discharged state, while the LNMO was partially charged, confirming again the limited lithiation mechanism. The LTO electrode showed a reversible capacity comparable to its initial capacity; the LNMO however showed a reduced capacity due to a significantly increased impedance in the reassembled half cell, which is attributable to the disassembly process. At all three temperatures, the loss of cyclable lithium occurs through the “limited lithiation” mechanism. The increased rate of capacity fade with increasing temperature indicates the efficiency of LTO should be more dependent on temperature in comparison to LNMO. However, the opposite behaviour was observed in back-to-back cells (Figure 6). When no cross-talk is allowed between the electrodes, increased temperature caused more severe reactions on LNMO compared to LTO. These ACS Paragon Plus Environment

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results show that the reactions on LTO which are initiated by cross-talk products are indeed more sensitive to temperature. It is observed from resistance measurements that the resistance of LTO does not change significantly on cycling, although there is a more pronounced increase at 55 ˚C. Even at 55 ˚C, the average resistance (see Figure S4) for the 50th cycle is ~13 Ω.cm2. For LNMO, there is a gradual increase in resistance over cycling which is again more significant at 55 ˚C. In all cases, the increased resistance does not become high enough to significantly reduce capacity, as evidenced by the clear end-points in the voltage profiles. Compared with the back-to-back cells, the initial resistances are smaller (~10 Ω.cm2 vs. ~20 Ω.cm2) but later reach higher values (~110 Ω.cm2) for the 55 ˚C sample (see Figure S3 and S4). The initially lower values are due to the absence of contributions from the lithium electrodes in the back-to-back cells. However, since capacity fading is slower in regular full cells than in back-toback cells, and cycling shifts to higher voltage windows, the cells are cycled at more harsh conditions overall, and combined with the effect of electrode interactions it is hence possible to observe higher resistances. Even though the resistance contribution to the capacity is minor, it is likely that it would start to play a major role in later stages near the cell break-down point.


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Figure 8. Individual voltage curves together with resistance measurements throughout the cycles of LNMO-LTO 3-electrode cells at room temperature (RT), 40 and 55 ˚C for (a) the first cycle, (b) the 10th cycle, (c) the 30th cycle and (d) the 50th cycle.

As mentioned in the introduction, gas formation associated with electrolyte oxidation and interaction of gas species between two electrodes has been shown to exist in other full cell chemistries.23,24,25,26 For instance, CO2 evolved after electrolyte oxidation can migrate and subsequently be reduced on the negative electrode.27 With an aim to gain further insight into the role of gas interaction in the cross-talk mechanisms for this system, a modified “back-to-back” cell using a bipolar electrode and contained within a single pouch has been tested (see experimental part for details). In this experiment, gas interactions are possible between the electrodes but interaction of species in the electrolyte is prevented. The cycling results of this cell is shown in Figure 9. This type of cell showed faster a slightly increased rate of capacity fade as compared to the regular back-to-back cell, which can be attributed to a slightly reduced coulombic efficiency (~0.97, compared to ~0.98 for the regular backto-back cell). As the voltage curves show, the cyclable lithium loss originates from the “limited delithiation” mechanism as in the regular back-to-back cells. This indicates that the cross-talk due to interaction of gas species is unlikely to significantly contribute to “limited lithiation” mechanism observed in regular full cells. Therefore, the capacity fade due to the “limited lithiation” mechanism is most likely primarily due to the interaction of dissolved species in the electrolyte with LTO, rather than gaseous species (e.g. CO2; nevertheless, the possibility of cross-talk due to the much slower process of gas evolution and subsequent re-dissolution cannot be completely ruled out). Such species can be soluble electrolyte decomposition products, for example, the oxidation of EC can generate radical cations on positive electrode.33 These reactive radical cations can participate in further side reactions in different parts of the cell, or be reduced on the negative electrode. Similarly, dissolved transition metal cations (Mn, Ni) can travel to the negative electrode and be reduced there.19 The reduction of such species on LTO could contribute to the observed fading mechanism. As stated in the introduction, fast fading of full cells with graphite electrode is ascribed to SEI damage caused by the incorporation of transition metal ions.12 Interaction of Mn ions can induce pores, cracks and increase the electronic conductivity of the graphite SEI.34 Mn in the SEI can coordinate with EC inside the SEI and facilitate its reduction35 or decrease the reduction barriers via an electrocatalytic mechanism.36 Therefore, migrated species could cause additional reduction of electrolyte rather than being reduced itself. One of the reasons for studying this full cell system was to eliminate this second possibility as LTO operates at a significantly higher potential than graphite and does not require a comparable SEI for its operation. However, it should still be noted that the nature of the interfacial reactions on LTO are still not well understood and the second possibility should not be neglected (i.e., possibility of catalytic effects). In any of these cases, at higher temperatures, either new reaction routes which consume more electrons at LTO become favorable, or the kinetics related to existing cross-talk reactions are accelerated. In order to get further insight into these mechanisms, we performed ex situ characterization of both electrodes obtained from regular full cells (two-electrode cells).


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Figure 9. Cycling performance of ‘single pouch’ back-to-back cells allowing only gas interactions (a) and the selected voltage curves from the same cell (b). Large symbols are capacity and small symbols are Coulombic efficiency.

Surface characterization of LNMO electrodes after 50 cycles. The morphology of the LNMO electrodes after cycling (50 cycles, C/5 rate) at room temperature and at 55 ˚C have been observed via SEM (Figure 10). Both conductive carbon and active LNMO particles showed no significant observable differences after cycling at either temperature. This indicates that there is no considerable damage to nor thick surface film formation on the LNMO particles. The same electrodes were also characterized via XPS to investigate the surface film thickness and composition.


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Figure 10. SEM images of pristine LNMO electrodes (a), after cycling at room temperature (b) and at 55 ˚C (c).

C 1s, O 1s and F 1s spectra of pristine and cycled LNMO electrodes obtained with the in-house XPS instrument (Al-Kα radiation 1486.6 eV) are shown in Figure 11a-c. The main differences are observed in the C 1s and O 1s spectra. For the pristine electrode, a large contribution appears from carbon black which is assigned to the peak at 284.5 eV (all spectra were calibrated with respect to hydrocarbon peak; 0.5 eV difference was fixed between the hydrocarbon and carbon black peaks). The green peaks in Figure 11 are assigned to the binder which is a co-polymer of PVdF and HFP (Kynar Flex 2801 binder, 88 mol% PVdF and 12 mol% HFP). In C 1s spectrum, C=O/O-C-O and C-O related adsorbed species could be observed before cycling. In the O 1s spectrum, a metal oxide peak (529.8 eV) is observed together with C=O and C-O related adsorbed species. The F 1s spectrum shows one large peak around 688 eV corresponding to the PVdF-HFP binder. After cycling, the binder-related peaks shift to lower energies, as previously reported.13,37 The main differences in both the C 1s and O 1s spectra are due to evolution of surface species which are related to organic surface film components as a result of electrolyte oxidation. For both temperatures, no significant differences can be observed when comparing the relative intensities of the surface species, except for a slight increase in the peaks for C=O related species for the 55 ˚C sample. Comparing the relative intensity of the peaks corresponding to the surface species to the peaks for the bulk in the C 1s spectrum reveals an increase in the relative amount of surface species with respect to the binder and CC peaks. No similar observation can be made in the O 1s spectrum in, for example, the relative ratio of surface species to the metal oxide (i.e. substrate) peak. This indicates that a comparatively larger deposition of surface species occurs on carbon black and binder-rich regions than on the active material when the temperature is 55 ˚C. In addition, slightly more LiF or M-F (e.g. MnF2, NiF2) can be seen at 55 ˚C. If general trends in atomic concentrations are compared (see Figure 11d), a significant decrease in F content is observed in accordance with intensity loss from the binder-related peaks. The decrease in concentration of C is comparatively smaller. Here, surface films increase C concentration while at the same time these films decrease the contribution from carbon black. The latter effect seems to dominate the overall trend in C concentration. For O 1s, the contribution from surface films dominates over the intensity loss from the bulk metal oxide supporting a slightly more preferential film formation on carbon black-binder regions. The spectra for P 2p, Mn 2p and Ni 2p are shown in Figure S6. ACS Paragon Plus Environment

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It is clear, and in agreement with the SEM observations, that there is no significantly thick surface film formed on LNMO. A thin layer is indeed observed, but its thickness does not seem to increase by cycling at elevated temperature as seen in the O 1s spectra. Since it is obvious that this layer does not stop electrolyte decomposition and also does not change in thickness, side-reaction products are not continuously deposited on active particle surface and instead released into the electrolyte or in a broad sense into the cell (e.g. gases). In the internal resistance measurements (Figure 8), it was seen that the major contributions to the overall resistance originated from the LNMO electrode. When the cells were opened for sample preparation, we observed deposited species on the separator facing the LNMO electrode which were partly adhered to the electrode outer surface. It is thus possible that highly reactive species formed after electrolyte oxidation react further with the polypropylene separator, creating deposits which partially block the pores of the separator.

Figure 11. C 1s, O 1s and F 1s XPS spectra of pristine and cycled LNMO electrodes.

Surface characterization of LTO electrodes after 50 cycles. As shown in Figure 12, and in contrast to the LNMO electrodes, the morphology of the LTO electrodes changes clearly on cycling. Surface films seem to be formed both on the LTO active particles and on the conductive carbon additives. Measurement of 100 carbon black particles gives an average particle size of 54 nm (±11 nm) in the pristine electrode (Figure 12a), which increases to 70 nm (±15 nm) after room temperature cycling and 76 nm (±13 nm) after 55 ˚C cycling. According to the SEM images of the LTO electrodes cycled at 55 ˚C, apparent surface layer formation can be observed since previous sharp edges and voids appear to be filled with previously unobserved deposits. According to the resistance measurements (Figure 8), relatively thicker layers could be expected for LNMO electrodes, but instead only an extremely thin layer exists on the surface according to SEM and XPS analysis. On the other hand, the LTO electrode did not contribute much to the overall internal resistance, but thicker SEI layer formation is indeed observed by SEM. We performed XPS analysis using both in-house XPS and high energy synchrotron XPS in order to gain further insight into the surface products deposited on the LTO electrodes. The synchotron photon energies of 2005 and 6015 eV probe to different depths within the surface which allows us to obtain depth information of the surface layers.


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Figure 12. SEM images of pristine LTO electrodes (a), after cycling at room temperature (b) and at 55 ˚C (c).

The survey spectra (see Figure 13) of the pristine sample display characteristic peaks of carbon (~285 eV), oxygen (~531 eV), and titanium (~459 eV). This is expected for an electrode comprising watersoluble binders and LTO. The most notable differences between the pristine and the cycled samples are the disappearance of the titanium peak, the emergence of the fluorine peak (~680 eV) and an increasing O/C ratio for the cycled samples. This is in accordance with previous results.38,31 The most notable difference is the increased F concentration in the 55 ˚C compared to the RT sample.

Figure 13. Survey spectra of pristine and cycled LTO samples measured using different photon energies.


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The C 1s spectra of the different samples are displayed in Figure 14 (note that the cycled samples were measured using three different photon energies). In the pristine electrode, the conductive carbon is identified as the peak at 284 eV in the pristine sample. General C-H bonds are found at 285 eV. Weak contributions to intensity are observed in the region where carboxylates (289.3 eV), O-C-O / C=O related species (288 eV) and ethers (286.7 eV) might be present.39 Since the exact ratio of binders used, occurrence of particle surface treatment or presence of additives are unknown, a detailed assignment cannot be made. The general trend in the C 1s spectra is that from the pristine to the cycled samples there is a significant increase in the C-O region, correlated to a decrease in C-C peak (which can be considered as a bulk peak). A peak at around 289 eV is also observed, corresponding to carboxylate groups. When comparing the different photon energies, for the room temperature (RT) sample large contributions are still visible at 284 eV, but is significantly decreasing with lower excitation energies. This indicates that the SEI is thick enough to barely allow the in-house XPS to probe through it. For the 55 ˚C sample, it is observed that no significant signal at 284 eV is observed at any photon energy, indicating an even thicker SEI at elevated temperatures. The relative intensity of organic species decreases as the photon energy increases (i.e. increasing probing depth) for both temperatures. A similar trend is also seen when the RT sample is compared with the 55 ˚C sample indicating that the main difference between them is due to film thickness rather than the film composition.

Figure 14. C 1s and O 1s spectra of the pristine, room temperature, and 55 ˚C LTO sample at different energies.

In the O 1s spectra, reference lines for C-O and C=O related species with binding energies of 533.5 eV and 532.2 eV were placed on the graph, respectively, while titanium oxide could be observed at ~530.5 eV.38,40 The different oxygen peaks in the organic compounds usually have a large overlap and are hard to separate, however, a lower binding energy can be expected for C-O contribution from ethers (532.7 eV) compared to C-O in carboxylates (*O-C=O, 533.6 eV).39 The most intense peak for the RT sample (at 532.7 eV) is in accordance with ethers, which is also observed in the C 1s spectrum. At 55 ˚C, the main peak becomes broader, indicating compositional changes in addition to the increase ACS Paragon Plus Environment

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of film thickness. This can be due to contributions from salt decomposition products (e.g. LixPFyOz) at 534.6 eV as well as C-O and C=O contributions from carboxylates. One interesting feature in the O 1s spectra is that the oxide peak, being clearly visible at 6015 eV excitation energy, hardly has any contribution at lower energies. The probing depth of the O 1s spectra can be estimated to be around 9, 14 and 40 nm for excitation energies of 1486.6, 2005 and 6015 eV, respectively.41 By comparison of the integrals of the metal oxide peak to the organic oxygen peaks, the SEI thickness can be assumed to be between 9-14 nm for the RT sample since the metal oxide peak becomes slightly visible when measured by photon energy of 2005 eV (blue curve). However, the SEI thickness is seen to be above 14 nm for the 55 ˚C sample (note that there is a number of assumptions made in estimations of these numbers and see SI for further discussion). As can be seen in Figure 15, the pristine sample contains neither phosphorous nor fluorine. Both phosphorous and fluorine are present in the hexafluorophosphate salt. The double peak feature in the P 2p spectra indicates that there are both P-O bonds and P-F bonds present. This suggests that salt decomposition products are incorporated into the SEI while some undecomposed salt may also be present. The F 1s spectra also shows signs of salt decomposition products, since the peak width indicates that several different components are present. One important difference is observed only for the RT sample in the region where contributions can come from LiF (or MnF2, NiF2). These species seem to be deeper in the SEI as they are only seen at high excitation energies. Absence of related peaks at 55 ˚C can be due to increased solubility of LiF or M-F species at this temperature, which could lead to partial dissolution of some SEI components. The spectra for Mn 2p, Ni 2p and Ti 2p are shown in Figure S7 and S8.

Figure 15. P 2p and F 1s spectra of the pristine, room temperature (RT), and 55 ˚C LTO sample measured using different photon energies.

In summary, it can be concluded that the cycled samples of LTO show signs of a relatively thick SEI comprising ethers, carboxylates, phosphates and F-containing species. Ni and Mn are obviously present in the SEI on the LTO anode (see SI). It is interesting to note that even though a considerably ACS Paragon Plus Environment

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thick surface layer is formed on the LTO surface, the impedance rise from LTO is still insignificant (Figure 8). Electrochemical testing of these cells showed that there should be steady side reactions occurring on the LTO side - meaning electrolyte oxidation products are able to reach the LTO surface through these layers. It is therefore not surprising that these layers do not cause considerable ionic resistance. These results also show the importance of individual internal cell resistance measurements combined with the XPS surface characterization, since it is common to correlate thick surface layers to cell resistance even when this is not supported by any electrochemical testing of cells.

XANES. The deposition of dissolved cathode transition metal ions on the anode surface is a wellknown example of ‘cross-talk’ in full cells (e.g. for systems with LMO, LNMO or NMC positive electrodes). The reduction of these migrating ions to their metallic state has been reported,11 however, more recent papers claim non-metallic oxidation states for Mn and Ni.22,42 In our previous study, Mn and Ni peaks were observed after only one cycle43 on LTO, and this has also been seen in this current work for both room temperature and high temperature cells. However, it was not possible to determine the oxidation state of the metals via XPS. Therefore, in order to determine these oxidation states, XANES measurements have been performed on cycled LTO electrodes as well as reference samples for different oxidation states of Mn and Ni. The results are shown in Figure 14 and indicate the Mn oxidation state to be close to +III and the Ni state to be close to +II. This shows migrated metal ions are not reduced to their metallic state, in agreement with the recent literature.44 Although metal ion deposition on the anode is not a direct cause for capacity loss via reduction to metallic state, it is still associated with the cross-talk and possibly with side reactions on LTO. It should be noted that temperature does not appear to have any effect on the oxidation state of the metal ions in the surface layer.

Figure 16. XANES measurements of LTO samples cycled at room temperature and 55 ˚C.

Conclusions In this study, a comprehensive analysis of capacity fading occurring in LNMO-LTO cells has been performed using different cell configurations and galvanostatic cycling combined with an intermittent ACS Paragon Plus Environment

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current interruption technique. One of the primary motivations in this study was also to understand the mechanisms causing the observed fast fading at elevated temperature. It is observed that severe parasitic reactions are occurring on the LNMO surface in both half cells and full cells. These reactions, in isolation, would be expected to cause fading of cells via a loss of cyclable lithium originating from an increasingly “limited delithiation” of both electrodes in the full cell. Electrode interactions (cross-talk) are in fact observed to result in a greater degree of parasitic reactions (i.e. a lower coulombic efficiency) at the LTO electrode relative to the LNMO electrode in the full cell. This lower efficiency of the negative electrode results instead in an increasingly “limited lithiation” of the electrode materials. Accelerated failure at elevated temperature results from significantly lower coulombic efficiency of LTO. Evolution of gases can also cause rapid cell failure, but can be mitigated with optimization of cell design. Stable capacity retention at room temperature may be misleading as an indication of cycle life in practice, as a significant degree of side reactions may occur even at room temperature, resulting in progressive decomposition of the electrolyte. Coulombic efficiency must therefore be taken into account, as it may be significantly below 100% even if there is no apparent fade in capacity. A typical lab-scale test cell such as those investigated in this work may cycle with good reversibility for more than three months, or 1000 cycles (see Figures S4 and S9) before exhibiting sudden capacity fade. This is partly because the electrolyte is in significant excess compared to realistic cells, and this variable should be considered when investigating long-term or high temperature cycling behaviour. According to the capacity fading mechanisms shown in this study, some strategies to increase the efficiency of the negative electrode (e.g., separator decorated with lithiated LTO or a larger particle size of LTO) may create a situation in which the efficiencies of both electrodes are equivalent and result in deceptively ‘stable’ cycling at certain temperatures. Based on our results, it might be expected that system with equivalent electrode efficiencies at high temperature may show a lower efficiency for LNMO at room temperature, which would be expected to result in the “limited delithiation” mechanism. Therefore, it is important to check the cycling stability at different temperatures and preferably over long durations to verify that such strategies do indeed improve the cell stability. However, based on the present results, we suggest that future strategies preferably target improving the efficiency of the LNMO electrode, such as oxidatively stable electrolytes, as this is ultimately the major source of inefficiency at LTO. Surface film formation after electrolyte oxidation and transition metal dissolution from the cathode are widely considered to cause failure in LNMO-based cells. However, resistance measurements and voltage curves obtained from three-electrode measurements confirm limited contributions to capacity loss from internal resistance (or overpotential increase) and active mass loss due to metal dissolution. The resistance increase in the system is observed to mainly originate from the LNMO side. The resistance increase was more pronounced at elevated temperature despite a lack of change in the thickness of the surface layer on LNMO, indicating a minor contribution to resistance from the surface layer. Resistance increase might instead result from precipitation of solids in the separator as indicated from our results, but may also result from other effects such as an increase in contact resistance between the electrode and current collector,45 or some reconstruction of the LNMO surface such as has been observed in layered oxide electrodes,46,47 which we have not investigated here. In any event, the resistance increase did not contribute significantly to capacity fade in our experiments. In contrast to the LNMO, surface films formed on the LTO electrode are found to be relatively thick but show comparatively little resistance. Those films also consisted of Mn and Ni migrated from LNMO and they are shown via XANES to be in a non-metallic state.

Supporting Information Supporting information includes the code used to numerically simulate the limited lithiation and delithiation mechanisms. Figures: Schematic view of lithiation states of LNMO and LTO under static aging conditions (S1). The change of self-discharge/cycle with cycle time (S2a) and voltage curves ACS Paragon Plus Environment

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(S2b) for the varying OCV experiment. The median resistance change with respect to LNMO (a) and LTO (b) electrodes in LNMO-LTO back-to-back cells (S3) and 3-electrode cells (S4). S5. Potential profiles of LNMO and LTO half cells from the cycled back to back (a) and 3-electrode cells (b). S6. XPS spectra of LNMO electrodes for P 2p, Mn 2p and Ni 2p. Discussion of thickness of surface films on LTO. XPS spectra of LTO for Mn 2p and Ni 2p (S7) and Ti 2p (S8). S9. Cycling results of LNMOLTO cell cycled for 2400 cycles. Acknowledgement We acknowledge Helmholtz Zentrum Berlin for synchrotron radiation beam time for HAXPES and XANES measurements. Dr. Maria Hahlin and Dr. Roberto Félix Duarte are acknowledged for their help during HAXPES and XANES measurements, Dr. William Brant for his help for data processing in Athena software. Leclanché are also gratefully acknowledged for providing the LTO electrodes. Volkswagen AG and Scania CV AB are acknowledged for financial support. The authors also acknowledge support from the Swedish Energy Agency (project no. 42031-1) and StandUp for Energy.

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Understanding the Capacity Loss in LNMO-LTO Lithium-Ion Cells at

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