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On the Anisotropic Lattice Strain and Mechanical Degradation of High- and Low-Nickel NCM Cathode Materials for Li-Ion Batteries Aleksandr Olegovi# Kondrakov, Alexander Schmidt, Jin Xu, Holger Geßwein, Reiner Mönig, Pascal Hartmann, Heino Sommer, Torsten Brezesinski, and Jürgen Janek J. Phys. Chem. C, Just Accepted Manuscript • DOI: 10.1021/acs.jpcc.6b12885 • Publication Date (Web): 17 Jan 2017 Downloaded from http://pubs.acs.org on January 21, 2017

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The Journal of Physical Chemistry

On the Anisotropic Lattice Strain and Mechanical Degradation of High- and Low-Nickel NCM Cathode Materials for Li-Ion Batteries Aleksandr O. Kondrakov,*,† Alexander Schmidt,† Jin Xu,*,‡ Holger Geßwein,‡,† Reiner Mönig,‡,† Pascal Hartmann,†,§ Heino Sommer,§ Torsten Brezesinski,*,† and Jürgen Janek†,# †

Battery and Electrochemistry Laboratory, Institute of Nanotechnology, Karlsruhe Institute of Technology, Hermann-von-Helmholtz-Platz 1, 76344 EggensteinLeopoldshafen, Germany. ‡

Institute for Applied Materials, Karlsruhe Institute of Technology, Hermann-vonHelmholtz-Platz 1, 76344 Eggenstein-Leopoldshafen, Germany. †

Helmholtz Institute Ulm for Electrochemical Energy Storage, Helmholtzstraße 11, 89081 Ulm, Germany. §

BASF SE, 67056 Ludwigshafen, Germany.

#

Institute of Physical Chemistry, Justus-Liebig-University Giessen, Heinrich-Buff-Ring 17, 35392 Giessen, Germany.

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Abstract In the near future, the targets for lithium-ion batteries concerning specific energy and cost can advantageously be met by introducing layered LiNixCoyMnzO2 (NCM) cathode materials with a high Ni content (x ≥ 0.6). Increasing the Ni content allows utilization of more lithium at a given cell voltage, thereby improving the specific capacity, but at the expense of cycle life. Here, the capacity fading mechanisms of both typical low-Ni NCM (x = 0.33, NCM111) and high-Ni NCM (x = 0.8, NCM811) cathodes are investigated and compared from crystallographic and microstructural viewpoints. In situ X-ray diffraction reveals that the unit cells undergo different volumetric changes of around 1.2 % and 5.1 % for NCM111 and NCM811, respectively, when cycled between 3.0 V and 4.3 V vs. Li/Li+. Volume changes for NCM811 are largest for x(Li) < 0.5 due to severe decrease of the interlayer lattice parameter c from 14.467(1) Å to 14.030(1) Å. In agreement, in situ light microscopy reveals that delithiation leads to different volume contractions of the secondary particles of (3.3 ± 2.4) % and (7.8 ± 1.5) % for NCM111 and NCM811, respectively. And post-mortem cross-sectional scanning electron microscopy analysis indicates more significant microcracking in the case of NCM811. Overall, the results establish that the accelerated aging of NCM811 is related to the disintegration of secondary particles caused by intergranular fracture, which is driven by mechanical stresses at the interfaces between the primary crystallites.

Introduction Layered lithium transition metal oxides of the type LiNixCoyMnzO2 (NCM, with x + y + z = 1) are well established and widely used cathode materials for rechargeable lithium-ion batteries (LIBs), mainly because of their excellent electrochemical properties.1,2 However, state-of-the-art applications place high demands on the cathode material(s), particularly when high specific energy is the primary concern. Thus, they need to be morphologically and structurally optimized so that the maximum performance in terms of capacity, stability, and safety is achieved within the electrochemical window of carbonate-based electrolytes (e.g., 3.0-4.3 V vs. Li/Li+).3 In the quasi-ternary phase system of NCM (LiNiO2/LiCoO2/LiMnO2), LiNiO2 has the highest gravimetric and volumetric capacity in this voltage range; consequently, increasing the nickel content in NCM is effective in improving the specific capacity, and therefore the energy density of full-cells with a graphite-based negative electrode.4-6 However, high-Ni NCM materials have lower structural and thermal stability, and typically show higher capacity fading.7 In recent years, significant efforts have been made to understand the mechanisms leading to capacity fading, gassing, and impedance rise in cells using Ni-rich NCMs.8-15 2 ACS Paragon Plus Environment

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It has been suggested that particularly Ni4+ species are highly reactive and tend to oxidize the electrolyte solvent(s), thus leading to CO2 evolution and solid electrolyte interphase (SEI) layer formation at the cathode.8-10,14,15 In addition, the formation of electrically insulating NiO at the NCM surface during cycling is believed to impair the electrode performance.16 Recent studies indicate that the capacity fading is also largely related to material fracture (mechanical degradation). Such degradation can occur within secondary particles because of significant lattice changes upon charging and discharging.17,18 In situ X-ray diffraction (XRD) performed on LiNi0.8Co0.1Mn0.1O2 (NCM811) showed that the lattice parameters exhibit a non-monotonic dependence on lithium content, and that the lattice changes become strongly anisotropic at high states of charge (SOC).18 In addition, it has been reported that the NCM811 lattice experiences the most significant changes in the interlayer distance at voltages beyond 4.0 V vs. Li/Li+. Unfortunately, the existing crystallographic data are only moderately resolved, particularly with respect to lithium content, and no unit cell position (zcoordinate) for the oxygen atoms is available in the literature. Thus, the transition metal (TM)-oxygen (TM = Ni, Co, Mn) and lithium-oxygen bond lengths are not known with sufficient precision. In addition, there are no detailed quantitative data on the anisotropic peak broadening, which reflects crystallographic strain. In view of the anisotropic volume changes, the net strain in secondary particles is strongly dependent on the relative orientation of adjacent crystallites, and its determination by means of XRD is very difficult. To gain a better understanding of the capacity fading mechanism(s), a more thorough and detailed investigation of the crystallographic and volumetric changes of high-Ni NCM materials in practical electrodes is required. Here, we report a comparative study of two NCM materials with different Ni contents, namely NCM111 (33 at.% Ni) and NCM811 (80 at.% Ni), by XRD, light microscopy, and scanning electron microscopy (SEM). The diffractometer used allowed obtaining XRD patterns directly on hermetically sealed pouch cells and the excellent time resolution of the experiments made in situ tracking of unit cell volume, bond lengths, and crystallographic strain possible. Overall, we analyze the microstructural changes as a function of lithium content and correlate the results with data – from in situ light microscopy and post-mortem SEM analysis – on the evolution of the secondary particle volume and mechanical degradation of the active material.

Experimental Section Cell Assembly and Electrochemical Cycling NCM-based positive electrodes, with a loading of approx. 2 mAh/cm2, were obtained from BASF SE and consisted of 93 wt. % active material. Information on packing 3 ACS Paragon Plus Environment


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density as well as on particle morphology and size distribution in the electrodes can be found in Fig. S1. 1 M LiPF6 in ethylene carbonate and ethyl methyl carbonate (3:7 by weight, LP57) was used as the electrolyte. The capacity retention tests were performed on 2032-type coin cells with 600 µm-thick lithium metal foil (Rockwood Lithium Inc.) and Celgard 2500 polypropylene separator using 100 µl of electrolyte. Prior to assembly, all cell compartments as well as the electrodes and separator were dried at 60 °C under vacuum. The cells were assembled in an Ar-filled glovebox or in a dry-room with dew point below −50 °C, and they were cycled at 45 °C and at C/2 using a MACCOR Series 4000 cycler (Tulsa). X-Ray Diffraction XRD patterns were collected using a laboratory diffractometer with Mo-Kα1,2 radiation designed for battery research. Details of the diffractometer setup can be found elsewhere.19 The high-flux beam (up to 108 photons per second) allowed obtaining in situ data on pouch cells cycled at 25 °C and at a rate of C/10 without using Kapton windows. A single XRD pattern was obtained during a time period of 75 s. Every two patterns were compared to eliminate noise arising from cosmic radiation and then merged. The sample-to-detector distance, detector tilt angles, and instrumental resolution function were determined using an annealed CeO2 powder sample. Both the detector calibration and azimuthal integration of the 2D images were performed with the pyFAI software package.20 For structural refinement, TOPAS Academic V5 software was used. The peak shapes were described using a pseudo-Voigt Thompson-CoxHastings profile function. The background was subtracted using a Chebyshev polynomial function. The lattice parameters were refined first, then the position of the oxygen atoms (z-coordinate), and finally the microstructure parameters.21 The initial structural models were chosen based on the refinement of neutron diffraction data obtained for pristine NCM powder samples at the high-resolution powder diffractometer for thermal neutrons (HRPT) at the Swiss continuous spallation neutron source (SINQ), Paul Scherrer Institute (PSI), Villigen, Switzerland (see Fig. S2 for details of the measurement parameters).22 Light Microscopy and Cross-Sectional SEM Electrodes for use in the optical cell were cast on aluminum foil with holes – having diameters ranging from 50 µm to 100 µm – separated by a distance of about 200 µm. Lithium metal (99.9 %, Alfa Aesar) was used as the counter electrode. Both electrodes along with a glass fiber filter separator (GF/A type, Whatman) were mounted onto a polyethylene frame. This frame was then sealed in a special home-built cell, equipped with a borosilicate glass window. A schematic of the experimental setup used is shown in Fig. S3. In this setup, the cathode is turned so that the active material is on top (i.e., 4 ACS Paragon Plus Environment

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oriented towards the microscope to ensure high-resolution imaging). The resulting overvoltages were determined to be 130 mV and 40 mV for NCM111 and NCM811, respectively. Nevertheless, the volume changes as a function of lithium content could be studied reasonably well. All the optical images were taken with an Olympus BXFM microscope, equipped with a CCD camera, in bright field mode. It should be noted that the low optical transparency of the electrolyte reduced the optical resolution of the microscope. High- and low-magnification SEM micrographs were collected on a MERLIN from Carl Zeiss. SEM samples of pristine and cycled NCM cathodes were prepared by ion beam milling, through cross-sectional thinning using the Hitachi ion milling system IM4000.

Results and Discussion Electrochemical Performance Fig. 1 shows the cyclability of NCM111 and NCM811-based coin-type cells at 45 °C. The cells were galvanostatically charged to 4.3 V vs. Li/Li+ and then the voltage was kept constant for 1 h or until the current dropped below C/50 (from initially C/2). Metallic lithium was used as the negative electrode, which allowed for direct comparison of the electrochemical behavior of both NCM materials. This is due, in part, to the fact that lithium is – unlike graphite-based electrodes – not affected by Mn and Ni poisoning, and is also capable of compensating for lithium losses in side reactions during cycling. Overall, under the conditions used in this work, lithium continuously reacts with the electrolyte until the cells dry out.23-26 As a result, the long-term performance is significantly worse than for full-cells containing a “stable” graphite anode. However, here we only compare the cyclability during the first 80 cycles, where the loss of electrolyte at the anode can still be neglected to some extent. The voltage profiles (top panel in Fig. 1) exhibit subtle plateaus (on charge) at 3.8 V and 3.75 V for NCM111 and NCM811, respectively. As expected, the Ni-rich material resembles the electrochemical behavior of LiNiO2 more closely.27 For example, the change in slope around 4.2 V suggests coexistence of several phases.28,29 The NCM111 and NCM811 half-cells delivered discharge capacities of 157 mAh/g and 193 mAh/g in the initial cycle and 137 mAh/g and 132 mAh/g in the 80th cycle (middle panel in Fig. 1). The average capacity fade rate per cycle was significantly larger for NCM811 compared to NCM111 (0.40 % vs. 0.16 %). This observation confirms the lower stability of high-Ni NCMs, as reported previously.6,7 The first cycle irreversibility was found to be 12 % and 14 % for NCM111 and NCM811, respectively, which is in the range typically observed for NCM materials.10,30,31 However, the Coulombic efficiency (see Fig. S4) stabilized well above 99 % after the initial cycle. The bottom panel in Fig. 1 shows the evolution of the mean 5 ACS Paragon Plus Environment


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discharge voltage, which reflects a buildup of overvoltage (and consequently an increase in cell resistance) with cycling. As seen, the mean discharge voltage decreased in a relatively linear manner for both materials, though at a higher rate for NCM811. Nevertheless, it should be noted that the lithium anode and other cell components also contribute to the increase in cell resistance.23,25 Collectively, the data in Fig. 1 establish that NCM811 cells show accelerated aging or, in other words, lower cycling stability than low-Ni NCM, which agrees with previous reports.7,17

Figure 1. Cycling performance of NCM111 and NCM811 half-cells at 45 °C. (top) Typical voltage profiles at C/10. Only the first two cycles are shown for clarity. (middle) Capacity retention and (bottom) mean discharge voltage at C/2. 6 ACS Paragon Plus Environment

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In Situ X-Ray Diffraction For in situ XRD experiments, pouch cells 4 × 2 cm2 in size were used. They were charged from the open circuit voltage (3.09 V and 3.26 V for NCM111 and NCM811, respectively) to the cutoff voltage of 4.3 V. The latter voltage was kept constant for 1 h in every cycle. During this time, the current dropped below C/200 from initially C/10, thereby indicating that the charge cycle was completed. The NCM111 and NCM811 half-cells delivered first cycle discharge capacities of 149 mAh/g and 189 mAh/g, respectively, and showed good electrochemical performance, comparable to that of coin-type cells (see Fig. 1). In total, more than 700 diffraction patterns were collected during the course of two charge/discharge cycles. The data for a single pattern were recorded during 150 s, corresponding to a lithium extraction ∆x(Li) ≈ 0.003 per pattern. Fig. 2 shows a contour plot illustrating the evolution of the NCM111 reflections (see also Fig. S5 for contour plot of peak intensities of an NCM811 half-cell, along with representative in situ XRD patterns). The high intensity of the 111 and 200 Al reflections leads to a significant overlap with the 006, 012 and 104 reflections of NCM. The 220, 311, and 222 Al reflections are much less intense and can be well distinguished from those of NCM at higher angles. Structural refinement confirmed the samples to be of layered ܴ3ത݉ type, in agreement with literature reports.18,32 At first glance, the continuous and steady evolution of Bragg reflections of both NCM materials indicates no structural phase transition with cycling. Nevertheless, it is known that pure LixNiO2 (being one component of the quasi-ternary NCM phase system) undergoes a transition from rhombohedral to monoclinic symmetry at intermediate lithium content (0.75 < x(Li) < 0.5).29,33 This symmetry reduction is accompanied by splitting of some reflections, such as 110 and 113. The XRD patterns of NCM811 suggest such a phase transformation. However, because of the limited angular resolution of our in situ setup and severe overlap with Al reflections from the foil of the pouch, a clear distinction between rhombohedral and monoclinic NCM could not be made (see Fig. S6 for line broadening/peak splitting and Rietveld plots for XRD patterns fitted in ܴ3ത݉ and ‫ܥ‬2/݉). Thus, the rhombohedral phase was used as the structural model in the Rietveld analysis of diffraction data (note that changes among the rhombohedral and the monoclinic phase are minor).33 The degree of cationic disorder for the lithium and nickel (Li/Ni exchange) was determined from neutron diffraction data for pristine NCM111 and NCM811 powder samples (see Fig. S2 for observed, calculated, and difference neutron diffraction profiles). The slightly higher disorder found in the latter material (6.8 % vs. 4.3 %) agrees with literature reports on high-Ni NCMs.34,35



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Figure 2. Contour plot of peak intensities of an NCM111 half-cell, along with representative in situ XRD patterns and voltage profiles.

Fig. 3 shows the results from Rietveld refinement, including voltage profiles, lithium content (x(Li), calculated from cycling data), lattice parameters (a, c), TM-oxygen bond length (lTM-O), lithium-oxygen bond length (lLi-O), unit cell volume (V), and anisotropic microstrain parameter (S202). As seen, the lattice parameters are very similar for the pristine materials, but vary significantly with lithium content (in charged state at 4.3 V, NCM111 and NCM811 retain 45.6 % and 25.4 % of the original lithium content, respectively). However, these variations are reversible. We note that (1) the delayed change in lattice parameters observed for NCM811 during the initial charge is not yet fully understood, but seems to be correlated with both the higher irreversibility and the overvoltage in the first cycle; and that (2) the apparent offset in lattice parameters at the 8 ACS Paragon Plus Environment

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beginning of the second cycle (with respect to the first cycle) is consistent with that in lithium content. In the following, we only focus on the crystallographic changes in the first charge cycle. The lattice parameter a gradually decreases from 2.8569(1) Å to 2.8145(1) Å for NCM111 and from 2.8702(1) Å to 2.8127(1) Å for NCM811. Since nickel, cobalt, and manganese ions have similar X-ray scattering factors and are statistically distributed within the lattice, XRD does not distinguish between different metal ions. The term “TM” then corresponds to a mixture of these three ions and the measured distances are averaged. It should be noted that the ܴ3ത݉ symmetry allows for variation in the oxygen z-coordinate (zO), which affects the TM and lithium-oxygen bond lengths (lTM-O and lLi-O). As a result, the accuracy of both the lTM-O and lLi-O estimations is limited by the accuracy of zO refinement. Although oxygen interacts only weakly with X-rays and the zO values show larger scattering than those of a and c, changes in zO could be resolved well (see Fig. S7). Formally, at x(Li) = 1, the nickel ions are in the oxidation state 3+ (assuming Mn3+ and Co3+), and the effective lTM-O is 1.959 Å and 1.961 Å for NCM111 and NCM811, respectively (assuming low-spin TM state and r(Ni3+) = 0.70 Å, r(Co3+) = 0.685 Å, r(Mn3+) = 0.72 Å and r(O2−) = 1.26 Å).36 These values agree with our experimentally determined bond lengths of 1.962(2) Å and 1.963(2) Å. The value of lTM-O for NCM111 decreases almost linearly with time (from 1.962(2) Å to 1.892(3) Å). lTM-O refined for NCM811 also decreases linearly from 1.963(2) Å to 1.896(3) Å with charging to 4.0 V, but jumps back to 1.909(2) Å at the end of cycle. Since the lTM-O difference corresponding to this jump (0.013 Å) is only slightly larger than the refinement error (0.005 Å), we consider these changes as insignificant. The lLi-O vs. time curve for NCM111 follows a monotonic, near linear behavior, with the refined lLi-O value increasing from 2.105(3) Å to 2.176(4) Å. In contrast, the curve for NCM811 exhibits more complex behavior; lLi-O increases from 2.114(4) Å to 2.162(5) Å up to 4.0 V (x(Li) = 0.5) and then decreases significantly to 2.117(5) Å with further charging to 4.3 V. Fig. 3 also shows that during charge the lattice parameter c of NCM111 increases from initially 14.218(1) Å to 14.500(1) Å at x(Li) = 0.5. Further delithiation leads to a slight decrease to 14.480(1) Å. The c value of NCM811 passes a maximum (14.467(1) Å) at x(Li) ≈ 0.5. Delithiation of NCM811 below x(Li) = 0.5 leads to a large decrease in c to 14.030(1) Å, which is even lower than in the initial discharged state (14.195(1) Å). Overall, these data demonstrate very clearly that NCM811 undergoes more pronounced structural changes than NCM111, which is why we discuss them in more detail in the next section.



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Figure 3. Structural refinement results for (left) NCM111 and (right) NCM811 cathodes cycled against lithium metal anodes in pouch cells. The evolution the cell voltage (E), lithium content (x(Li)), lattice parameters (a, c), TM-oxygen bond length (lTM-O), lithiumoxygen bond length (lLi-O), unit cell volume (V), and anisotropic microstrain parameter (S202) is shown.

The values of a and lTM-O directly depend on the oxidation states and ionic radii of TMs in NCM811. The decrease in both values during charge reflects that lithium ions are removed from the lattice and the charge is compensated by oxidation of TMs. More drastic changes in the TM-oxygen bonding are therefore related to the higher redox state(s) (and higher specific capacity) compared to NCM111. The lithium-oxygen bond length plotted in Fig. 3 corresponds to an average distance between the lithium and the oxygen in LiO6 octahedra and depends on both a and c (or, more rigorously, also on zO). As seen, lLi-O rather follows the trend observed for c, thereby indicating that the expansion and contraction of the lattice along the c direction occurs predominantly in the lithium slabs (TMO6 octahedra contract nearly linearly during the charge cycle, as demonstrated by decreasing lTM-O). 10 ACS Paragon Plus Environment

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Although we do not fully understand the explicit nature of the non-monotonic changes in c and lLi-O, we can speculate about the cause based on crystal and electronic structure data from first principles for isomorph LiNiO2 and LiCoO2.37-40 The interlayer distance in layered lithium transition metal oxides has been shown to be controlled by the repulsion between the TM-oxygen slabs.39,41 Thus, the increase of the c value corresponds to an increase of the interslab repulsion and vice versa. Initially, a large number of lithium ions screen the negative charge of oxygen, thereby diminishing the interslab repulsion. Upon lithium extraction, the screening effect of the lithium layer gradually vanishes and the interslab repulsion increases, which ultimately leads to the lattice expansion along the caxis. At very high degrees of delithiation, the screening effect does not hinder the interslab repulsion anymore. Having this is mind, contraction of the lattice along the caxis during charge, as observed for NCM811, must imply a decrease in the interslab repulsion. Theoretical studies predict that the TM 3d and O 2p states are highly hybridized in layered high-Ni LiTMO2 compounds and form a strongly mixed TM-O band.38 This means that electrons can be extracted from both TM and O states. In such a case, the effective negative charge of the oxygen atoms – and therefore also the interslab repulsion – decreases, thereby causing shrinkage of the c-axis. The decreasing screening effect of lithium should facilitate the interslab repulsion, inducing elongation of the c-axis. The decreasing effective oxygen charge should reduce the interslab repulsion and cause shrinkage of the c-axis. The data show that both effects do not compensate each other completely. Overall, we believe that the screening effect determines the interlayer distance in the range 1 > x(Li) > 0.5, while below x(Li) = 0.5 it seems to be controlled mainly by the effective charge of the oxygen atoms. A possible explanation for this may be that the mixing of TM and O states does not depend linearly on the lithium concentration. Thus, this would mean that O states do not participate in the charge compensation right from the beginning of delithiation, but rather beyond x(Li) = 0.5, where the TM-O bands start to overlap. The changes in unit cell volume reflect the changes that active material crystallites undergo during the charge/discharge process. Large volume changes inevitably generate significant mechanical strain inside the particles, which may lead to fracture and pulverization, and thus to formation of reactive surfaces and undesired side reactions with the electrolyte. For this reason it is important to analyze the unit cell volume as a function of SOC to determine potential voltage limits for reliable battery operation or, in other words, to find the optimum between specific energy and structural stability. Fig. 3 shows the unit cell volume vs. the time from in situ XRD; changes in lithium content calculated on the basis of the integrated current are shown as well. During charge, NCM111 and NCM811 undergo volume shrinkage from 100.50(1) Å3 to 99.33(1) Å3 and 101.27(1) to 96.19(1) Å3, respectively, which agrees with data available in the literature.17 The changes in lattice parameter a during charge are rather small 11 ACS Paragon Plus Environment


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(0.07(1) Å), but affect the unit cell volume by the second power. This is well visible at the beginning of the charge cycle. While a and c change in opposite directions, the unit cell volume resembles the evolution of a. However, above 4.15 V (x(Li) = 0.5) and 4.04 V (x(Li) = 0.5) for NCM111 and NCM811, respectively, the variations in lattice parameter a are less significant, whereas c still exhibits significant changes. Consequently, the unit cell volume changes at high voltages are determined by the lattice parameter c. We note that the specific capacity delivered by the NCM811 cathode above 4.04 V is only around 47 mAh/g, corresponding to 25 % of the total capacity; however, the volume change is quite significant (3.5 % from the total volume change of 5.1 %). Since large changes have a detrimental effect on the stability of the active material, the upper voltage limit is critical for the cycling performance. Both the faster capacity decay and higher rate of impedance rise observed for NCM811 compared to NCM111 (see Fig. 1) may be caused by the volume changes occurring at voltages between 4.04 V (x(Li) = 0.5) and 4.3 V (x(Li) = 0.25), a hypothesis that still requires experimental proof. Besides the net volume changes, it is also important to consider microstrain effects during cycling. The microstrain can be estimated by analyzing the XRD line broadening.21,42,43 The diffraction patterns in Fig. 2 (see also Figs. S5 and S6) clearly indicate such line broadening, which varies non-monotonically with the diffraction angle. The broadening is most significant for the 10l and 01l reflections (i.e., for 015, 107, and 018). Croguennec et al. demonstrated that the unusual broadening of these reflections can be explained by statistically distributed stacking faults in the layered NCM structure.29 Stacking faults are generated if the parent O3-type structure with characteristic ABCABC oxygen stacking is locally broken into a structure of O1-type with AB oxygen stacking. Analysis of stacking faults is very challenging for the large number of diffraction patterns collected. Thus, we analyzed the anisotropic line broadening using the phenomenological model by Stephens.21 For the ܴ3ത݉ space group (hexagonal setting), there are four independent microstrain parameters SHKL: S400, S004, S202, and S301.44 Simultaneous refinement of all four microstrain parameters during the sequential Rietveld analysis did not yield stable results. As a result, only S202 was used, as this parameter had the largest effect on the fit quality. Fig. 3 shows the evolution of the microstrain parameter S202 for NCM111 and NCM811 with cycling. This reversible increase in S202 with delithiation is well reproduced in two cycles. The NCM111 S202 value increases up to E = 4.15 V (x(Li) = 0.5) and then levels off, whereas the S202 profile of NCM811 exhibits a clear maximum at E = 4.04 V (x(Li) = 0.5). It should be noted that the Stephens model does not allow figuring out the physical reasons for the crystallographic microstrain, but it quantifies homogeneity of the lattice parameters within the material. The S202 profiles indicate that the lattice changes occur inhomogeneously at high degrees of delithiation, likely due to scattering of the lattice 12 ACS Paragon Plus Environment

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parameters (a, c) among individual crystallites. A direct comparison of the S202 profiles of NCM111 and NCM811 reveals that the latter material is subjected to more severe, periodic buildup of microstrain during single charge/discharge cycles.

In Situ Light Microscopy and Cross-Sectional SEM In situ light microscopy (LM) was used to measure the size of the secondary particles of both NCM111 and NCM811 during charge and discharge. Fig. 4a shows a representative image taken through the optical window of the cell (see also SEM micrograph in Fig. 4b for comparison). As seen, the primary crystallites of the secondary particles can be resolved reasonably well. The relative change in secondary particle size throughout cycling was determined by digital image correlation (DIC).45 Large irreversible variations in size were observed during the course of the first cycle; the changes in later cycles were more reversible. In XRD, the first and following cycles showed the same behavior. From this, we conclude that effects such as surface layer (cathodic SEI) formation contribute to the data in the initial cycle in the in situ LM experiments. Thus, we used data from the second cycle for comparison with the XRD results. The volume changes were calculated from the size measurements, as obtained from DIC (see Appendix S1 for more details).

Figure 4. Microscopy of NCM811 cathodes. (a) Representative in situ light microscopy image. The primary crystallites are clearly visible. (b) SEM micrograph showing spherical particles with micrometer-scale dimensions.

Fig. 5 shows the evolution of volume of the secondary particles with time, along with E and x(Li). Comparing the in situ LM data with the results from XRD shows that there is a direct correlation between the changes in secondary particle volume and those in unit cell volume. As expected, NCM111 experiences smaller variations than NCM811. The NCM111 secondary particles shrink reversibly by (3.3 ± 2.4) vol.% upon delithiation, 13 ACS Paragon Plus Environment


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whereas those of NCM811 undergo a (7.8 ± 1.5) vol.% shrinkage during charge. The larger data spread for NCM111 is likely due to the partially porous and less spherical morphology, resulting in non-radial or anisotropic movement of grain boundaries within single particles. In contrast, the secondary particles of NCM811 are nearly perfectly spherical in shape and have a rather dense and homogeneous structure. It seems that they follow more closely the crystallographic changes of the primary crystallites, and therefore the data are less spread out (lower standard deviation). The volume changes of NCM811 are most pronounced in the voltage range from 4.0 V to 4.3 V. The profile in Fig. 5 shows a strong decrease in secondary particle volume above 4.0 V, which agrees well with the shrinkage of the unit cell volume, as determined by Rietveld refinement (see Fig. 3).

Figure 5. Changes in secondary particle volume from in situ light microscopy for (left) NCM111 and (right) NCM811 cathodes cycled against metallic lithium anodes. Only the second cycle is shown for clarity. Shaded areas represent the error margin. The corresponding voltage profiles and the evolution of lithium content (x(Li)) with cycling time are shown as well.

Fig. 6 presents cross-sectional SEM micrographs taken from NCM cathodes before and after cycling at C/2. As mentioned above, the pristine NCM111 particles have a porous microstructure with visible grain boundaries between the primary crystallites, whereas 14 ACS Paragon Plus Environment

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the NCM811 particles reveal a more densely packed morphology. There are some microcracks in the cycled NCM111, which, however, cannot be distinguished from the pore structure initially present in the material. In contrast, the microcracks observed for NCM811 can be clearly attributed to the electrochemical cycling, as such defects are not present in the pristine material. The SEM micrograph of NCM811 shows a uniformly spaced network of multiple microcracks. We note that similar crack patterns have been reported for high-Ni NMCs before.17,46 The initially “dense” particles release mechanical stresses that build up during cycling by forming microcracks along the crystallite boundaries. Their zigzag shape is indicative of intergranular fracture. The width of the microcracks decreases towards the particle surface, thereby indicating that the formation begins in the interior, likely due to the presence of tensile stresses. This kind of stress typically arises if the regions in the interior of particles contract more than those at the surface, or if the surface expands more than the regions closer to the center. The first scenario is not very realistic, however, because the particle interior is not likely to be delithiated prior to the surface. The second scenario seems to be more plausible. During discharge, the near-surface regions expand more rapidly than those in the interior, given that only the top surface of the NCM particles is in direct contact with the electrolyte and conductive carbon black additive. It should be noted that the assumption of spherical lithium concentration profiles is a simplification. Considering a strong dependence of the electrical conductivity on crystal orientation in layered lithium transition metal oxides, the shape of the lithium concentration profile can be affected significantly by the texture of the secondary particles.47 Accordingly, it is reasonable to assume that complex lithium concentration gradients exist within the polycrystalline NCM particles, which in turn can lead to local stresses and strains. A recent study by means of transmission X-ray microscopy showed that lithium concentration profiles have irregular, non-spherical shapes.48 Interestingly, these profiles bear some resemblance to the crack patterns observed in the present work.

Figure 6. Cross-sectional SEM micrographs at different magnifications of (a-d) NCM111 and (e-h) NCM811 cathodes collected before (a, b, e, f) and after (c, d, g, h) 50 cycles at C/2. 15 ACS Paragon Plus Environment


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The strong anisotropic lattice changes observed by XRD may be responsible for the intergranular fracture. Grain boundaries of different types exist when the secondary particles are made of randomly distributed (oriented) primary crystallites. Small-angle boundaries will not experience large stresses since the adjacent crystallites expand in a similar manner. In contrast, boundaries with large angles (e.g., those, where {100} and {001} facets meet) will expand inhomogeneously and large stresses are expected to occur. In this context, it should be noted that local stresses may induce material fracture at the grain boundaries. The microstrain parameter S202 determined for NCM111 and NCM811 indicates that the crystalline lattice is significantly strained at high voltages. This may lead to imperfections such as stacking faults, and their accumulation eventually to mechanical degradation.42,43 Active material fracture is inevitably accompanied by the formation of (new) reactive surfaces and detrimental side reactions with the electrolyte. In addition, it may also lead to loss of electrical contact of active electrode regions, all of which helps explain the cycling data shown above.

Conclusions In summary, structural and morphological changes of NCM111 and NCM811 upon electrochemical cycling have been investigated by means of in situ X-ray diffraction, in situ light microscopy, and scanning electron microscopy. Cycling tests at 45 °C showed that the difference in stability is, among others, related to the impedance buildup in the cathode. X-ray diffraction revealed that NCM811 exhibits more significant (reversible) changes in lattice parameters than NCM111, when cycled in the voltage range between 3.0 V and 4.3 V with respect to Li/Li+. Refinement of the transition metal-oxygen bond length suggested that variations in bond character, similar to that reported for LiNiO2, may be responsible for the structural changes of NCM811 during cycling. Light microscopy confirmed that NCM811, in fact, undergoes much larger volume contraction upon charge, with the most pronounced changes occurring at E > 4.0 V, in agreement with the result from X-ray diffraction. As expected, based on these data, the Ni-rich material showed severe microcracking. Collectively, our research data suggest that the accelerated aging (i.e., faster capacity fading and higher rate of resistance buildup) of cells using NCM811 is closely related to the intrinsic degradation of high-Ni NCM materials caused by significant anisotropic lattice changes – particularly the abrupt and non-monotonic changes in c-axis – at high states of charge. To ensure a reliable performance of high-Ni NCM materials in LIBs, we need to identify (structural and chemical) ways that allow for compensation of mechanical stress and reduction of surface reactivity. 16 ACS Paragon Plus Environment

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Author Information Corresponding Authors *Phone: +49 721 60826870, E-mail: [email protected] *Phone: +49 721 60822679, E-mail: [email protected] *Phone: +49 721 60828827, E-mail: [email protected]

Notes The authors declare no competing financial interest.

Associated Content The Supporting Information is available free of charge on the ACS Publications website. Low-magnification cross-sectional SEM micrographs of NCM111 and NCM811 cathodes; Rietveld plots for neutron diffraction patterns for NCM111 and NCM811; schematic of the in situ light microscopy setup; Coulombic efficiency data for NCM111 and NCM811; in situ XRD patterns and Rietveld plots for NCM811; evolution of the oxygen z-coordinate in NCM111 and NCM811; and details on the calculation of the relative volume changes of secondary particles (Appendix S1).

Acknowledgements This study is part of the projects being funded within the BASF International Network for Batteries and Electrochemistry. We gratefully acknowledge the assistance of Dr. Denis Sheptyakov provided during the neutron diffraction measurements at the HRPT beamline of the Swiss continuous spallation neutron source SINQ, Paul Scherrer Institute, Villigen, Switzerland.

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and Low-Nickel NCM Cathode Materials for Li-Ion ... - ACS Publications

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