Investigation into Mechanical Degradation and Fatigue of

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Investigation into Mechanical Degradation and Fatigue of HighNi NCM Cathode Material: A Long-Term Cycling Study of Full Cells Simon Schweidler, Lea de Biasi, Grecia Garcia, Andrey Mazilkin, Pascal Hartmann, Torsten Brezesinski, and Jürgen Janek ACS Appl. Energy Mater., Just Accepted Manuscript • DOI: 10.1021/acsaem.9b01354 • Publication Date (Web): 05 Sep 2019 Downloaded from pubs.acs.org on September 6, 2019

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ACS Applied Energy Materials

Investigation into Mechanical Degradation and Fatigue of High-Ni NCM Cathode Material: A Long-Term Cycling Study of Full Cells Simon Schweidler,† Lea de Biasi,†,‡ Grecia Garcia,† Andrey Mazilkin,†,§ Pascal Hartmann,†,# Torsten Brezesinski,*,† and Jürgen Janek*,†,∥

† Battery

and Electrochemistry Laboratory, Institute of Nanotechnology, Karlsruhe Institute of

Technology (KIT), Hermann-von-Helmholtz-Platz 1, 76344 Eggenstein-Leopoldshafen, Germany. ‡

Institute for Applied Materials - Energy Storage Systems, Karlsruhe Institute of Technology

(KIT), Hermann-von-Helmholtz-Platz 1, 76344 Eggenstein-Leopoldshafen, Germany. §

Institute of Solid State Physics, Russian Academy of Sciences, Ac. Ossipyan str. 2, 142432

Chernogolovka, Russia. # BASF ∥

SE, Carl-Bosch-Strasse 38, 67056 Ludwigshafen, Germany.

Institute of Physical Chemistry & Center for Materials Science, Justus-Liebig-University

Giessen, Heinrich-Buff-Ring 17, 35392 Giessen, Germany.

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ACS Applied Energy Materials 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Abstract Nickel-rich NCMs [Li1+x(Ni1−y−zCoyMnz)1−xO2] are among the most promising cathode materials for use in high-energy lithium-ion batteries. While the cathode composition can vary depending on the application, graphite remains the material of choice at the anode side. In this article, we study the degradation of Li1+x(Ni0.85Co0.1Mn0.05)1−xO2 (NCM851005) in practical graphite-based full cells over hundreds of cycles by a combination of operando X-ray diffraction, electrochemical impedance spectroscopy, and advanced electron microscopy, and correlate the results with data from galvanostatic charge/discharge measurements at 1C rate and 45 °C. In addition, half-cells were assembled from the cycled positive and negative electrodes for better understanding of the fatigue behavior. Electrochemical testing revealed virtually linear capacity decay and impedance growth with cycling, which can be attributed predominantly to NCM851005 degradation. Although electron microscopy indicated that severe fracture of the NCM851005 secondary particles occurred, the capacity fading is found to be due not only to mechanical degradation and/or side reactions, but also to continuous surface reconstruction on the primary particle level from layered to rock salt-like structure, thus building up a kinetic barrier.

Keywords lithium-ion battery, lithium nickel cobalt manganese oxide, graphite, chemo-mechanical degradation, interfacial fracture, surface reconstruction

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Introduction Layered lithium transition metal oxides of type Li1+x(Ni1−y−zCoyMnz)1−xO2 (NCM or NMC) are among the state of the art cathode materials for Li-ion battery (LIB) applications. By increasing the Ni fraction, the specific capacity for a given cutoff voltage can be increased, with Mn exerting a stabilizing effect on the lattice structure and Co reducing cation mixing in the Li layers.1–3 However, increasing the specific capacity or, in other words, the Ni fraction in the NCM typically adversely affects the cell cyclability and longevity. In general, Ni oxidation during charge is accompanied by generation of highly reactive Ni4+ species, leading to electrolyte decomposition at the NCM particle’s surface. This, in turn, leads to gas evolution and formation of a cathode solid electrolyte interphase (cSEI) layer, thereby increasing the internal cell pressure and impedance. Furthermore, it has been shown by Lin et al. and Huang et al. that especially highly delithiated NCM is unstable; oxygen loss and migration of transition metal ions into the Li layer lead to formation of spinel and/or rock salt-like phases.2,4,5 Likewise, recent studies have demonstrated that there is a correlation between the capacity fading and the chemo-mechanical degradation of NCM secondary particles. With increasing Ni fraction, the primary particles suffer from more and more severe volume changes during (de-)lithiation, eventually leading to interfacial fracture and mechanical failure with cycling operation.6–11 In this work, we report a comprehensive study on the long-term cycling performance of singlelayer pouch cells using NCM851005 (85% Ni) and graphite as cathode and anode, respectively. Specifically, electrochemical results obtained on both full cells and half-cells are correlated with (micro-)structural data from operando X-ray diffraction and electron microscopy to better understand the degradation and fatigue behavior.

Experimental Section Materials and Electrochemical Testing NCM851005 cathodes of loading around 2 mAh/cm2 and with 94 wt% active material were obtained from BASF SE. They were calendared to a density of 3.2 mg/cm3. LP472 (1M LiPF6 in a 3:7 mixture [wt:wt] of ethylene carbonate and diethyl carbonate and 2 wt% vinylene carbonate) was used as electrolyte (BASF SE). Graphite anodes of loading around 7.1 mg/cm2 and with 96 wt% active material were also obtained from BASF SE. Specific capacities are given based on the total mass of active material (graphite or NCM851005) in the electrode. Electrochemical (long-term) testing was performed on single-layer pouch-type full cells. They were assembled inside a dry room with dew point below −60 °C and comprised a NCM851005 ACS Paragon Plus Environment


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cathode (50·50 mm2), a Celgard 2500 polypropylene separator (53·53 mm2), a graphite anode (51·51 mm2), and 500 µl of LP472 electrolyte. The full cells were first subject to a formation cycle, involving charging at C/10 (1C = 192 mA/g) to 3.7 V (25 °C), followed by a resting period of 24 h at constant voltage (CV) and 45 °C. Then, they were transferred to a dry room, cut open to release evolved gasses, and vacuum sealed again. Subsequently, the full cells were subject to a rate performance test at 25 °C (Figure S1), and eventually, they were cycled at 1C and 45 °C in the voltage range of 2.8-4.2 V. At the end of each charge cycle, a 1 h CV step at 4.2 V was implemented. All cells were cycled in Binder KB 720-UL climate chambers using a MACCOR

Series

4000

battery

cycler.

More

than

20

independent

cells

were

charged/discharged, as described above, and stopped after 100, 150, 200, 250, 300, 500, and 750 cycles for further characterization. After reaching the desired number of cycles, they were cycled at C/10 and 25 °C for three more cycles. In the third cycle, a direct current internal resistance (DCIR) measurement was performed (more details in Figure S2). As to structural characterization, the full cells were analyzed by operando X-ray diffraction (XRD). Subsequently, they were disassembled in a glovebox, and the electrodes were rinsed with 3 ml anhydrous dimethyl carbonate (≥99 %; Sigma-Aldrich). After drying, electrodes (Ø14 mm) were punched from the cycled cathodes and anodes and reassembled into CR2032 coin-type half-cells with a Li metal counter electrode (Ø15 mm; Albemarle Germany GmbH) using a GF/A glass microfiber separator (Ø17 mm; Whatman) and 200 µl LP472 electrolyte. Half-cells (three independent cells for each cycle number) with either NCM851005 or graphite electrode were then cycled at C/10 and 25 °C in the voltage range between 2.9 and 4.3 V and 0.005 and 1.2 V vs. Li+/Li, respectively, with a 30 min CV step at the upper cutoff voltage. Electrochemical Impedance Spectroscopy Electrochemical impedance spectroscopy (EIS) was performed on a Bio-Logic VMP-3 potentiostat (Bio-Logic SAS). A three-electrode Swagelok cell setup was used for the experiments, allowing individual evaluation of both electrode impedances.12–14 To this end, working/counter electrodes (Ø12 mm) were punched from the cycled full cell cathodes and anodes. The reference electrode was prepared by pressing Li metal into the 0.5 mm reference bore holder, and it was positioned in plane with the working/counter electrodes. GF/A (Ø13 mm) and LP472 were used as separator and electrolyte, respectively. EIS was performed at 50% state-of-charge (SOC) in the frequency range between 200 kHz and 10 mHz. The threeelectrode full cells were cycled at C/10 and 25 °C in the voltage range of 2.8-4.2 V. X-Ray Diffraction


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A laboratory diffractometer with a molybdenum micro-focus rotating anode (Mo-Kα1.2) and a fast area detector (Pilatus 300 K-W) was used for operando XRD. XRD was conducted on the full cells after electrochemical testing and the corresponding DCIR measurement. More details about the setup used and calibration procedure can be found elsewhere.6,15,16 During XRD, the cells were cycled at C/10 and 25 °C in the voltage range between 2.8 and 4.2 V, with a 1 h CV step at the upper cutoff voltage. 2D patterns were recorded in transmission geometry with an exposure time of 90 s in the 2θ range of 7.1-44.3°. Generated spike noise from cosmic radiation was corrected for by comparing and adding up the intensity of two consecutive patterns, resulting in an effective total time resolution of 180 s. Structural refinement was done using TOPAS Academic V5 software. Electron Microscopy Scanning electron microscopy (SEM) images at different magnifications of the pristine and cycled NCM851005 and graphite electrodes were taken on a LEO-1530 electron microscope (Carl Zeiss AG). For cross-sectional SEM, samples were prepared by ion-beam polishing through cutting using an IB-19510CP (Jeol Germany GmbH). Transmission electron microscopy (TEM) was performed on a Titan 80-300 (FEI Company) Cs image corrected transmission electron microscope operated at 300 kV. Samples for TEM investigation were prepared using a STRATA dual-beam focused ion beam (FIB) system (FEI Company). The FIB lift-out samples were milled using a Ga-ion beam at 30 kV, followed by final milling at 2 kV to improve the surface quality. Scanning TEM (STEM) images were collected using a highangle annular dark-field (HAADF) detector. Electron energy-loss spectroscopy (EELS) data were acquired with an energy resolution of 1 eV (Tridium 863 imaging filter, Gatan Inc.), as determined by measuring the full-width at half-maximum of the zero-loss peak. The energy positions of the transition-metal L edges were calibrated using the zero-loss peak.

Results and Discussion Electrochemical Testing of Pouch-Type Full Cells Figure 1a shows the evolution of voltage profiles over 750 cycles of a graphite/NCM851005 pouch cell at 1C rate and 45 °C (after formation and subsequent rate performance testing). The charge/discharge curves are characteristic of graphite-based cells using high-Ni NCM.6,17 The initial specific discharge capacity is around 195 mAh/gNCM. As expected, capacity fading is observed during cycling, with losses of 6, 15, 22, and 28% after 200, 400, 600, and 750 cycles, respectively (Figure 1b). Of note, the fluctuations in specific capacity and Coulombic ACS Paragon Plus Environment


efficiency (every 50 cycles; Figure 1a and b) are due to implemented resting periods of varying lengths. The data scattering seen after around 380 cycles is not fully understood yet, but may be related to partial blocking of the free electrode surface by gas bubbles.18 In the literature, the capacity fading of full cells is frequently attributed to various factors, such as lithium loss at the anode side.19,20 Apart from that, the cathode is also prone to degradation due to anisotropic lattice changes during (de-)lithiation, among others.7,8 In this regard, it is assumed that especially mechanical degradation leads to particle disintegration associated with loss of particle connectivity, increasing the cell impedance.21,22 With increasing cycle number, the charge curve shifts toward higher voltages, as can be seen in Figure 1a. This leads to shortening of the voltage plateau, initially present in the range of 4.1-4.2 V. In addition, the specific capacity achieved in the 1 h CV step at the upper cutoff voltage is found to increase steadily. This continuous shift/truncation of the voltage plateau becomes clearer by considering the differential capacity plots shown in Figure 1c. Note that the same behavior is also observed in half-cells (see section on electrochemical testing of halfcells below). The dq/dE curves demonstrate that the NCM851005 undergoes a series of phase transitions upon delithiation, from the hexagonal H1 via monoclinic M to the hexagonal H2 and H3 phases.22–24 These phase transformations are similar to those found for LiNiO2 (LNO), which is the Ni-containing end member of the NCM family.23,25–29 In the case of LNO, the H2/H3 phase transition region is accompanied by a sudden contraction of the c lattice parameter, leading to secondary particle fracture and capacity degradation.25 Consequently, it can be assumed that NCM851005 also suffers from structural instability due to anisotropic volume changes, imparting stress into the material.7 The dq/dE peak denoting the H2/H3 phase transition decreases gradually in size with increasing cycle number and vanishes after around 300 cycles. Likewise, all other phase transformation regions shift toward higher voltages. Nevertheless, the dq/dE curves indicate the reappearance of the H2/H3 (or H3/H2) phase transition in the discharge cycles, though at different overpotentials. This result thus demonstrates that, due to kinetic limitations, the high-voltage phase transformation only occurs during the course of the 1 h CV step at the end of charge (for cycle numbers >300).21,30–32


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Figure 1. (a) Voltage profiles of a graphite/NCM851005 pouch cell at 1C and 45 °C in the range between 2.8 and 4.2 V. (b) Specific charge/discharge capacity and Coulombic efficiency as a function of cycle number. (c) Differential capacity plots.

The measured DCIR of graphite/NCM851005 cells at 50% SOC is shown in Figure 2. As can be seen, it increases linearly (with increasing scattering) from about 100 Ω cm2 after 100 cycles to 380 Ω cm2 after 750 cycles. Note that both cathode and anode affect the cell impedance to different extents. An increase in resistance at the cathode side may be due to structural, morphological, and/or chemical changes of the NCM851005.7,23 At the anode side, for example, ACS Paragon Plus Environment


deposition of transition metal species onto the graphite is known to poison the SEI layer, thereby also increasing the cell resistance.33 For that reason, coin-type half-cells using cycled graphite and NCM851005 electrodes, harvested from the full cells, were characterized electrochemically to elucidate whether the cathode or the anode is predominantly responsible for the performance decay.

Figure 2. DCIR measurements conducted on graphite/NCM851005 pouch cells at 50% SOC after different cycle numbers.

Electrochemical Testing of Coin-Type Half-Cells and 3-Electrode Full Cells Figure 3a and b depicts the 2nd cycle specific charge/discharge capacity of pristine and fatigued NCM851005 and graphite half-cells at C/10 rate and 25 °C. As expected, the NCM851005 shows significant capacity loss with increasing cycle number. For unknown reason, the initial specific capacity is slightly lower than that achieved after 100 cycles. Nevertheless, the capacity decreases continuously from around 200 mAh/gNCM (for pristine NCM851005) to 140 mAh/gNCM (after 750 cycles). This behavior is also reflected in the corresponding dq/dE curves (Figure 3c). In the charge cycles, distinct shifts in peak location toward higher voltages and peak broadening are observed, with the H2/H3 peak vanishing after around 500 cycles. However, reverse phase transformation occurs during discharge, similar to the full cell observations. It therefore appears that the long-term cycled NCM851005 can still reach the H2/H3 phase transition region if sufficient time is allowed during the CV step at the upper cutoff voltage. Also, in line with expectations, the specific capacity gained in the CV period increases as the cycle number increases (Figure S3). This result suggests that (de-)lithiation is still possible, but kinetically hindered due to formation of a surface (blocking) layer and/or migration of transition metals into interstitial sites. The former is probably accelerated by mechanical degradation of the cathode material.


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In contrast to the NCM851005, graphite showed little capacity fading, which is also evident from the dq/dE curves (Figure 3d). There are no significant changes in the differential capacity plots with increasing cycle number, and the characteristic peaks of the different lithium intercalation stages can be readily identified. Even after 750 cycles, the cycling behavior is still reminiscent of that of pristine material. Taken together, we conclude that graphite has a negligible effect on the full cell capacity fading. This also indicates that poisoning of the anode only plays a secondary role in the performance degradation of graphite/NCM851005 cells.

Figure 3. Second cycle specific charge/discharge capacity of half-cells using (a) NCM851005 or (b) graphite and (c and d) the corresponding differential capacity plots. The coin cells were cycled at C/10 and 25 °C in the voltage range of 2.9-4.3 V and 0.005-1.2 V vs. Li+/Li, respectively. The cycle number indicates the number of cycles the electrode material experienced prior to harvesting from the pouch cell.

For a more detailed characterization of the resistance built-up during cycling operation, EIS measurements were conducted at 50% SOC on both pristine/fatigued NCM851005 and graphite electrodes using three-electrode cells (Nyquist plots in Figure 4a and b). Compared with the pristine cathode, the impedance of the positive electrode increased significantly after 750 cycles (to >800 Ω cm2). According to the model by Levi et al.,34 the charge-transfer resistance ACS Paragon Plus Environment


(semicircle at medium frequencies) represents the major contribution to the impedance. In contrast, the EIS spectra of the graphite anode remain virtually unaltered during cycling (minor impedance increase by about 20 Ω cm2 after 750 cycles). Note that the experimental data were collected ex situ using different anodes and cathodes harvested from pouch cells. Thus, minor inconsistencies in the evolution of impedance spectra may arise from electrode inhomogeneities and some variation among the cells tested. Nevertheless, the DCIR and EIS results reveal the same trend and demonstrate that the increase in cell impedance can be attributed predominantly to the NCM851005.

Figure 4. Nyquist plots of pristine and fatigued (a) NCM851005 and (b) graphite electrodes at 50% SOC. The data in (a) and (b) are offset by 40 and 8 Ω cm2 along the imaginary Z axis, respectively, for clarity. The cycle number indicates the number of cycles the electrode material experienced prior to harvesting from the pouch cell.

Structural Characterization Operando XRD was conducted on the graphite/NCM851005 pouch cells after 1 (1st cycle after cell formation and rate performance testing), 100, 150, 200, 250, 300, 500, and 750 cycles at 1C rate and 45 °C to gain more insight into the capacity fading mechanism. Rietveld analysis


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was performed assuming α-NaFeO2-type structure with R3m space group for NCM851005 and hexagonal symmetry P63/mmc for graphite. The P63/mmc crystal structure was applied to the full range of charge and discharge, although it is known that graphite undergoes a transformation to P6/mmm at higher SOCs.16,35–38 Because of superposition of the graphite reflections with high-intensity reflections of Al from both the cathode current collector and the pouch bag and Cu from the anode current collector, only the 002 reflection of graphite was clearly detectable. Nevertheless, the 002 reflection yields the interlayer spacing between the graphene sheets, which is a good measure of the degree of lithiation.16 Likewise, the 003 reflection of NCM851005 is well suited for monitoring the (de-)lithiation process of NCM-type cathode materials.7,39 Figure 5a and b shows representative XRD patterns taken before the charge cycle as well as after charge and discharge. The Rietveld refinement results (i.e., lattice parameters, unit cell volume, and phase fractions) can be found in Figure S4 and Table S1. Lattice parameter changes of NCM cathode materials as a function of SOC have already been described in considerable detail in the literature.6–8,39,40 In general, the c-axis increases in the beginning of charge, and therefore, the 00l reflections shift toward smaller Bragg angles. Then, at relatively high SOCs, the c-axis collapses, leading to a reverse shift of the 00l reflections toward larger 2θ values. The evolution of c lattice parameter with (de-)lithiation varies with the NCM composition or, in other words, the Ni content in the material. After 1 and 100 cycles, “complete” shift of the 003 reflection to 2θ = 8.8° at the end of charge is observed (Figure 5a). As the number of cycles increases, this shift becomes smaller, indicating progressively smaller contraction of the lattice along the c direction. Additionally, broadening of the 003 reflection is observed. Such broadening is evident after 150 cycles and, according to Rietveld analysis, it is due to coexistence of two NCM851005 phases of slightly different lattice parameters (XRD data could be refined only reasonably well when considering two phases).23 Note that the difference in lattice parameters is indicative of different degrees of delithiation. As mentioned previously, for LNO, phase transformation from the hexagonal H2 to hexagonal H3 phase occurs around 4.1 V. In that case, splitting of reflections is clearly observed, allowing unambiguous assignment of the H2 and H3 phases. Here, the two NCM phases can be distinguished by their different lattice parameters. From the above dq/dE curves, we conclude that the NCM851005 indeed undergoes a phase transformation, similar to LNO. However, the material rather shows solid-solution behavior, and there are resolution limitations of the diffractometer used, which certainly somewhat contributes to reflection broadening. In the following, we refer to the NCM851005 with the 003 reflection at relatively larger and smaller



Bragg angles as more active and less active, respectively (for pouch cells cycled between 150 and 500 cycles).

Figure 5. (a) Evolution of the 003 reflection of NCM851005 and (b) 002 reflection of graphite during cycling of graphite/NCM851005 pouch cells at C/10 and 25 °C. The blue, pale blue, and dark blue patterns were taken before starting the charge cycle (under open circuit voltage condition), in charged state (including 1 h CV step at 4.2 V), and in discharged state (2.8 V), respectively.

Figure 6a and b shows the refined c lattice parameter and estimated SOC for the less and more active NCM851005 phases at the end of charge as a function of cycle number (see Figure S5 for details about the determination of SOC from the c lattice parameter).41 With increasing cycle number, the fraction of more active NCM851005 decreases, while that of the less active phase increases until only the latter remains after 750 cycles. This result suggests complete loss of the more active NCM851005 with prolonged cycling. In addition, it can be seen that the c lattice parameter of both phases undergoes smaller changes as the cycle number increases, which is also reflected in the steadily decreasing SOC. After 750 cycles, about 74% SOC was determined by c lattice parameter comparison. This value is in good agreement with the SOC (71%) calculated from the electrochemical data. Similar behavior is also observed for the graphite anode. Clear splitting of the 002 reflection at the end of charge (Li intercalation into graphite) is seen up to cycle number 300, indicating ACS Paragon Plus Environment

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transformation of the Li-poorer LiC12 phase (11.6°) into the Li-richer LiC6 phase (11.0°) (Figure 5b). The LiC6 phase fraction decreases continuously with increasing cycle number, and after 500 cycles, only the LiC12 phase is formed upon cell charging. The disappearance of both the LiC6 (stage 1) reflection and the 003 reflection of the more delithiated (active) NCM851005 phase correlates with one another. Overall, it seems that the cathode experiences severe material fatigue during long-term cycling, such that the NCM851005 is no longer able to achieve high SOCs. Hence, the amount of lithium extracted at the cathode side eventually becomes insufficient to enable full lithiation of graphite (LiC6 formation).

Figure 6. (a) c lattice parameter and (b) SOC of the less and more active NCM851005 phases at the end of charge (including 1 h CV step at 4.2 V) versus the cycle number. The gray shaded area represents the cycle numbers where two NCM851005 phases were considered in the Rietveld analysis.

Figure 7 shows the c lattice parameter together with the charge trace as a function of specific capacity after 1, 250, and 750 cycles. Note that only the 1st and 750th cycle XRD data could be refined reasonably well assuming a single NCM851005 phase. In the 1st cycle, the expected ACS Paragon Plus Environment


characteristic c lattice parameter evolution is observed.6 The specific capacity gained during the 1 h CV step at the end of charge is negligible and so are the lattice parameter changes, indicating that the cell has attained “full” equilibration under the given conditions. In contrast, after 250 and 750 cycles, the specific capacity increased by 11 and 13 mAh/gNCM, respectively, in the CV step, and thus, the c lattice parameter decreased further. These results emphasize the increasing kinetic limitations of the fatigued NCM851005 cathode. Finally, a graphite/NCM851005 cell after 250 cycles at 1C rate and 45 °C was analyzed in some more detail by operando XRD. Specifically, it was charged at C/10 rate to 4.2 V, followed by a 60 h CV step. Figure S6 shows the voltage curve and the corresponding current profile as well as the evolution of the 003 reflection of NCM851005 and of the 002 (for LiC12) and 001 (for LiC6) reflections of graphite over time. After about 2 h CV charging, convergence of the two overlapping reflections denoting the less and more active NCM851005 phases is observed, accompanied by a continuous shift toward larger Bragg angles. After 9-10 h, XRD indicates “full” equilibration. Within this time, the specific capacity increased by about 20 mAh/gNCM, and the 003 reflection approached 2θ = 8.8°, similar to the cell after formation and rate performance testing (note that comparable specific charge capacities were achieved with 186 vs. 188 mAh/gNCM; Table S1). This result thus indicates that the fatigued NCM851005 can still be (almost) fully delithiated and the overall process is indeed kinetically limited.

Figure 7. Effect of the CV step at the end of charge on the c lattice parameter of NCM851005. Note that only refinement results for the 1st, 250th, and 750th cycles at C/10 and 25 °C are shown ACS Paragon Plus Environment

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together with the corresponding charge trace. Significant reflection broadening is observed after 250 cycles, and thus, two NCM851005 phases were considered in the Rietveld analysis.

Morphological and Microstructural Changes Chemo-mechanical degradation associated with interparticle contact loss is frequently reported to explain the capacity decay of NCM-based LIB cells upon cycling.21,30–32 In the following section on electron microscopy, we show that mechanical degradation is only indirectly responsible though, by fostering surface reconstruction or, in other words, formation of a resistive surface reduced layer, thereby increasing the cell impedance. SEM was used to examine the morphological changes on the primary and secondary particle levels. Figure 8a and b shows top view and cross-sectional images of the pristine NCM851005 cathode. From Figure 8a, the presence of some fractured secondary particles is evident, which we believe is due to electrode calendaring prior to cycling. Figure 8b indicates that the particles are rather dense, with slightly increasing porosity toward the core. After 100 and 500 cycles (Figure 8c-f), more fractured secondary particles are observed, some of them having a flattened structure at the top surface. However, they show different fracture patterns compared to the pristine NCM851005, and it is reasonable to assume that they suffered from cracking due to electrochemical cycling. Figure 8d and f indicate that microcracks are generated along grain boundaries and propagation takes place in a zigzag fashion (see also cross-sectional SEM images in Figure S7). The width of these cracks apparently decreases from the particle core to the surface, indicating that their formation begins in the interior. This result is in agreement with data available in the literature.7,22,23,40,42 Such cracking eventually creates channel-type structures, leading to side reactions of the freshly exposed (reactive) NCM851005 surfaces with the electrolyte, and consequently, to cSEI formation which, in turn, adversely affects the lithium transport through the electrode (across interfaces). Unlike NCM851005, the graphite anode barely changes from a morphological point of view during the course of cycling (Figure S8). The only noticeable change is that the initially smooth surface is being covered by a coarse but relatively uniform SEI layer. For that reason, we specifically focus on the cathode hereafter.



Figure 8. Top view and cross-sectional SEM images of the NCM851005 cathode in the discharged state (a and b) before cycling and (c and d) after 100 and (e and f) 500 cycles at 1C and 45 °C.

To better understand the mechanism(s) leading to capacity fading, the microstructural changes of NCM851005 were thoroughly investigated by STEM and TEM. Figure 9a-d shows highresolution STEM (HR-STEM) images obtained on the pristine and cycled material. They reveal the layered to rock salt-like phase transformation in a thin layer at the particle surface of the cycled NCM851005. The bulk of the primary particles shows alternating bright (transition metal ions) and dark atomic planes (Li ions), reflecting the layered lattice structure. However, the contrast in the surface regions is clearly altered due to site exchange defects (migration of transition metal ions into Li sites). Fast Fourier transformation (FFT) patterns confirm the phase transformation. Note that the presence of rock salt-like surface phases has been reported to have a profound effect on the cell cyclability by increasing the charge-transfer resistance, among others.23,43–45 While pristine NCM851005 shows no such detrimental surface layer, the thickness of the rock salt-like phase is found to increase from about 2 nm after 100 cycles to 14 nm after 500 cycles. However, the surface layer on the individual primary particles seems not uniform throughout. This is particularly evident for the material after 500 cycles (Figure S9). Phase transformation was only clearly observed for surfaces that apparently were exposed to the electrolyte (due to NCM851005 secondary particle fracture during cycling). In contrast, “intact” interface regions between adjacent primary particles did not show signs of rock saltlike phase formation.


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Figure 9. HR-STEM images of NCM851005 in discharged state (a) before cycling and (b) after 100, (c) 250, and (d) 500 cycles at 1C and 45 °C and the corresponding FFT patterns from the different particle regions. The rock salt-like surface layer is denoted by yellow lines.

These results were also corroborated by TEM (Figure 10a-c). HR-TEM of the interior of a fractured NCM851005 secondary particle after 250 cycles shows that the microstructure changes from the primary particle surface to the core. FFT from the bulk of the particle reveals a single crystalline pattern, while the reflections in the intermediate region become more diffuse and exhibit some splitting. Furthermore, small miss-oriented grains are found in the outer region of the particle, a result that might be explained by continuous impairment of the rock salt-like surface layer due to lattice oxygen release and gradual phase transformation.22,24,46–49 Unfortunately, the FFT results do not allow for clear conclusions about the crystal structure of the primary grain, as the patterns can be indexed to both layered and rock salt-like phases. Thus, EELS analysis was performed on the same areas. EELS spectra for the Mn, Co, and Ni L edges are shown in Figure S10. As somewhat expected, chemical shifts of the L3 and L2 edges to lower energy loss are found for the areas closer to the primary particle surface. The largest shift is observed for the Mn L edge, with ~2.7 eV for Mn L3 (~1.4 and 0.5 eV for Co L3 and Ni L3/L2 edges, respectively). Gradual decrease in the L3/L2 edge intensity ratio from the bulk to the surface is also seen for all three transition metal ions. Both the chemical shifts and the behavior of the L3/L2 edge intensity ratio are consistent with literature reports for Mn, Co, and



Ni in different valence states.50–52 Overall, EELS provides direct evidence of cation reduction (Mn+4 to Mn+2 and Co and Ni from +3 to +2) close to the particle surface. Taken together, we conclude from the electron microscopy data that there is a clear change from the layered bulk structure via an intermediate layer (possibly a mixture of layered, rock salt-like, and spinel-type phases) to the outer polycrystalline rock salt-like phase.

Figure 10. TEM of the interior of an NCM851005 secondary particle after 250 cycles at 1C and 45 °C. (a) Low-magnification TEM images showing secondary particle fracture, the red dashed squares denoting the area used for HR-TEM (b). Colored rectangles indicate the areas from which FFT patterns (c) were extracted: particle edge (blue), intermediate region (green), and particle core (red). The FFT from area 1 can be indexed to layered 𝑅3𝑚 structure, with zone axis along the [441] direction. In the FFT pattern from area 3, only the indices for the face-centered cubic rock salt structure are shown. The other reflections partially belong to the layered phase.

Conclusions In the present work, the degradation of practical graphite/NCM851005 cells was examined by galvanostatic charge/discharge tests, electrochemical impedance spectroscopy, X-ray diffraction, and electron microscopy. Electrochemical testing was performed at 1C rate and 45 °C in the voltage range of 2.8-4.2 V and revealed continuous capacity fading from initially around 195 to 140 mAh/gNCM after 750 cycles. Both impedance spectroscopy and direct current internal resistance measurements showed that the overall cell resistance increases with cycling, the origin of this impedance build-up being mainly the charge-transfer resistance of the cathode. ACS Paragon Plus Environment

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X-ray diffraction revealed the presence of two NCM851005 phases of slightly different lattice parameters during charge after about 150 cycles. However, fatigued NCM851005 (250 cycles) was found to be still capable of achieving almost full equilibration if sufficient time was allowed at the upper cutoff voltage, thus indicating kinetic limitations. As opposed to graphite, scanning electron microscopy demonstrated increasing fracture of the NCM851005 secondary particles with increasing cycle number, the microcracks being generated along grain boundaries in a kind of zigzag fashion. High-resolution (scanning) transmission electron microscopy further revealed continuous surface reconstruction on the primary particle level (for electrolyteaccessible surfaces). In conclusion, we have shown that the capacity fading is in fact associated to some degree with the mechanical degradation of the cathode material, even though indirectly. Anisotropic volume changes during (de-)lithiation impart stress into the material, leading to particle fracture, and therefore, inter-particle contact loss and adverse side reactions such as cathode solid electrolyte interphase layer formation. However, the freshly exposed and reactive NCM851005 surfaces seem very prone to phase transformation from layered to rock salt-like structure, all of which strongly contributes to impedance growth and performance degradation.

Associated Content Supporting Information Additional data from electrochemical testing of full cells and half-cells; details about DCIR measurements and SOC determination by comparison of c lattice parameters; Rietveld refinement results; and electron microscopy data.

Author Information Corresponding Authors *E-mail: [email protected]; phone: +49 721 60828827. *E-mail: [email protected]; phone: +49 641 9934500.

ORCID Simon Schweidler: 0000-0003-4675-1072



Lea de Biasi: 0000-0001-8546-0388 Grecia Garcia: 0000-0002-2675-2089 Andrey Mazilkin: 0000-0001-7805-8526 Torsten Brezesinski: 0000-0002-4336-263X Jürgen Janek: 0000-0002-9221-4756

Notes The authors declare no competing financial interest.

Acknowledgements This study is part of the projects being funded within the BASF International Network for Batteries and Electrochemistry and this work was partly carried out with the support of the Karlsruhe Nano Micro Facility (KNMF, www.knmf.kit.edu), a Helmholtz Research Infrastructure at Karlsruhe Institute of Technology (KIT, www.kit.edu). The authors thank Dr. Holger Geßwein and Dr. Reiner Mönig for providing access to the XRD.

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