Article pubs.acs.org/cm
Investigation of Changes in the Surface Structure of LixNi0.8Co0.15Al0.05O2 Cathode Materials Induced by the Initial Charge Sooyeon Hwang,†,‡,§ Wonyoung Chang,*,† Seung Min Kim,⊥ Dong Su,‡ Dong Hyun Kim,† Jeong Yong Lee,§ Kyung Yoon Chung,† and Eric A. Stach*,‡ †
Center for Energy Convergence, Korea Institute of Science and Technology, Seoul 136-791, Republic of Korea Center for Functional Nanomaterials, Brookhaven National Laboratory, Upton, New York 11973, United States § Dept. of Materials Science and Engineering, KAIST and Center for Nanomaterials and Chemical Reactions, IBS, Daejeon 305-701, Republic of Korea ⊥ Carbon Convergence Materials Research Center, Korea Institute of Science and Technology, Wanju-gun 565-905, Republic of Korea ‡
S Supporting Information *
ABSTRACT: We use transmission electron microscopy (TEM) to investigate the evolution of the surface structure of LixNi0.8Co0.15Al0.05O2 cathode materials (NCA) as a function of the extent of first charge at room temperature using a combination of high-resolution electron microscopy (HREM) imaging, selected area electron diffraction (SAED), and electron energy loss spectroscopy (EELS). It was found that the surface changes from the layered structure (space group R3̅m) to the disordered spinel structure (Fd3̅m), and eventually to the rock-salt structure (Fm3̅m), and that these changes are more substantial as the extent of charge increases. EELS indicates that these crystal structure changes are also accompanied by significant changes in the electronic structure, which are consistent with delithiation leading to both a reduction of the Ni and an increase in the effective electron density of oxygen. This leads to a charge imbalance, which results in the formation of oxygen vacancies and the development of surface porosity. The degree of local surface structure change differs among particles, likely due to kinetic factors that are manifested with changes in particle size. These results demonstrate that TEM, when coupled with EELS, can provide detailed information about the crystallographic and electronic structure changes that occur at the surface of these materials during delithiation. This information is of critical importance for obtaining a complete understanding of the mechanisms by which both degradation and thermal runaway initiate in these electrode materials.
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INTRODUCTION Due to their widespread application in portable electronic devices, lithium ion batteries (LIBs) are one of the most important technologies in everyday life. Because they employ the lightest and most electropositive metal for charge transfer (lithium), they have superior energy density,1 up to five times higher than that of lead acid batteries. Because of this high energy density, LIBs are being considered as a candidate for large-scale applications such as electric vehicles (EVs).2 There are, however, significant hurdles that must be overcome in order for them to see widespread usage in automotive applications. In contrast with mobile devices where lifetime requirements are only a few yearsbatteries in EVs are expected to be in use for more than a decade. This places substantial demands on their performance, specifically with respect to degradation in their energy storage capacity. In addition, EV applications require large batteries to achieve extended range: this leads to significant safety concerns, not only because larger batteries have a greater risk of explosion but also because the explosions are inherently more massive when they do occur. © 2014 American Chemical Society
In specific, Ni-based layered cathode materialsLiNi0.8Co0.15Al0.05O2, conventionally referred to as NCAare considered to be promising materials for EV applications because of their high discharge capacity (∼200 mAhg−1).3 However, these materials exhibit both a drastic capacity fade and an impedance rise with cycling, as well as poor thermal stability. Accelerated cycling tests at elevated temperature indicate that NCA cathode materials undergo significant power fading, which results in the development of an electrochemically inactive NiO-like phase with the rock-salt structure at the surface.4,5 It has also been shown that severely charged (i.e., “highly delithieated” or “overcharged”) Ni-rich cathode materials form a complex structure that is composed of a layered R3̅m core, a spinel structure subsurface, and a rock-salt structure at the surface.6 It is also well known that the original layered R3m ̅ structure of NCA readily becomes Received: October 10, 2013 Revised: December 22, 2013 Published: January 3, 2014 1084
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(0.1≤ x ≤0.5) using TEM and EELS. The outstanding spatial resolution of TEM allows us to distinguish subtle changes in crystal structure at the surface after electrochemical delithiation. EEL spectra are used for analyzing the electronic structure of charged NCA; we have focused particularly on the oxygen K-edge, because oxygen evolution is a crucial issue to battery safety. Additionally, because of the fact that XAS is usually obtained at higher energies, the O K-edge is generally less frequently investigated than the edges of the transition metals in these materials. We show that the degree of evolution in local surface structures differs among particles charged to the same potential, indicating that kinetic effects play a role in determining the overall structure. We have performed a statistical analysis utilizing EEL spectra to gain a deeper understanding of how the surface changes as a function of the state of charge. This study indicates that even “minor”, local changes in the structure and chemistry of NCA surfaces can have important impact on the safety characteristics of these materials, as these changes are the points of initiation for both battery degradation and thermal runaway.
thermally unstable when in the charged state, which results in phase transitions from the layered R3̅m structure to the disordered spinel structure (Fd3̅m) and finally to the rock-salt structure (Fm3̅m) with increasing temperature.7 These phase transitions are accompanied by the release of oxygen, which can accelerate thermal runaway by reacting with the flammable electrolyte inside the battery, thereby leading to catastrophic explosions.8,9 These studies clearly indicate that the substantial safety issues associated with NCA cathode materials are closely related to the existence of a structural instability. As a result, there have been a number of X-ray diffraction (XRD) studies that have investigated the evolution of the average crystallographic structure of the cathode materials as a function of either temperature or degree of delithiation.6,7,10,11 Our recent study of LixNi0.8Co0.15Al0.05O2 using in situ, time-resolved X-ray diffraction (TR-XRD) coupled with mass spectroscopy (MS) has suggested the possibility that surface degradation occurs in the half-charged state.11 Upon heating, oxygen gas was detected, coincident with the first phase transition of the layered structure to the spinel structure; this occurred even though a change in the oxygen stoichiometry is not required during this transition. It was presumed that some portion of the particle was more charged than the average charge state, thereby making that local area thermally unstable. However, because X-ray diffraction is a global technique, it was hard to prove where or how much of this area was contributing to the unexpected oxygen evolution that occurred upon heating. Thus, even though in situ TR-XRD/MS was found to be a novel tool to correlate structural changes and gas evolution, a complementary method is required to elucidate how these phenomena occur at the nanoscale. Phase transitions occur in layered transition metal oxides through the reduction and subsequent rearrangement of the transition metal ions12 accompanied by changes in the electronic density of states within the material. In prior work, Yoon et al. utilized X-ray absorption spectroscopy (XAS) to show that a change in electronic structure occurs in NCA with depth of charge.13 XAS, which offers excellent energy resolution, is well suited to the examination of the average electronic structure of electrode materials. However, XAS generally has poor spatial resolution, making it difficult to analyze a material with any site-specificity. In addition, because of the difference in the experimental set-ups required for each technique (e.g., vacuum levels, optics), it is difficult to integrate XRD and XAS together to allow simultaneous study of crystallographic and electronic structure changes with these techniques. Electrochemical charging or discharging is the process of inserting or extracting Li ions by reaching a designated potential value for the whole electrode. Therefore, analytical tools that collect signals from the bulk can play a significant role in understanding how cathode or anode materials behave during charging or discharging. However, both the degradation of electrode materials and the initiation of thermal runaway can sometimes start very locally within electrode materials. Transmission electron microscopy (TEM) is an ideal tool to investigate local structural changes and can provide information that is complementary to XRD and XAS. When combined with electron energy loss spectroscopy (EELS), TEM allows us to gain information concerning both the crystallographic and electronic structures of a material, as well as the morphology of a local area at high spatial resolution. Here, we report the local evolution of the surface structure of LixNi0.8Co0.15Al0.05O2 (NCA) as a function of state of charge
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EXPERIMENTAL SECTION
Electrochemical Testing. Positive electrodes (LixNi0.8Co0.15Al0.05O2; commercial product) were electrochemically delithiated to a half-charged (x = 0.5) and overcharged (x = 0.1) state at a rate of C/10 using a galvanostatic condition (i.e., constant current). The cathode was prepared as a mixed slurry of 90 wt % of active electrode material, 6 wt % of conducting Denka black, and 4 wt % PVDF binder in a n-methyl pyrrolidone (NMP) solvent. The mixed slurry was coated onto an Al foil that acted as a current collector. The 2032-type of coin cells were assembled with Li metal for an anode, a Celgard separator, and an electrolyte of 1 M LiPF6 dissolved in ethylene carbonate (EC), ethyl methyl carbonate (EMC), and dimethyl carbonate (DMC) solvent (1:1:1 by volume). The amount of lithium remaining in the cathode (x in LixNi0.8Co0.15Al0.05O2) was estimated from the charge and mass of the active material, with an assumption of 100% Coulombic efficiency. Transmission Electron Microscopy. The charged positive electrodes were thoroughly washed with a pure DMC solution to remove residual salts then gently abraded from the Al foil to acquire particles for the TEM experiment. Powder samples in a small vial of pure DMC solution were sonicated to ensure that they were well dispersed before dropping the solution onto a lacey carbon TEM grid. Sample preparation and loading into the TEM sample holder were done in an argon-filled glovebox, and the sample holder was transferred from a hermetically sealed container to the microscope in 10 s (at most) to minimize the exposure of the sample to air and moisture. Bright field (BF) and high-resolution electron microscopy (HREM) images and selected area diffraction patterns (SADPs) were obtained at room temperature with a JEM-2100F (JEOL) and a Tecnai G2 F20 (FEI), both operating at an accelerating voltage of 200 kV. High angle annular dark field (HAADF) images of overcharged NCA were acquired with a Titan 80−300 (FEI) equipped with a probe Cs corrector and a monochromator, operating at an accelerating voltage of 300 kV. Electron Energy Loss Spectroscopy. EEL spectra of the oxygen K-edge were acquired from pristine, half-charged (x = 0.5), and overcharged (x = 0.1) NCA using the JEM-2100F (JEOL) equipped with a GIF Tridiem spectrometer (Gatan). The spectra were obtained in diffraction mode with the selected area (SA) aperture defining the area that was analyzed. Spectra were aligned with reference to the zero loss peak position. The background of all spectra was subtracted using the power law method embedded within Digital Micrograph (Gatan) software. The energy resolution determined by a zero-loss peak was about 1.0 eV. 1085
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transformation (FFT) from the region indicated in Figure 2a results in a pattern that is identical to the SADP obtained from the bulk; therefore, we conclude that crystal structure of the NCA particle is homogeneous from the bulk to surface before electrochemical reaction. This was found to be the case for all pristine NCA particles analyzed. After 50% of the lithium is extracted from the pristine state, we can observe a structural modification at the exterior of the particles. Figure 3a shows a surface region of a half-charged
Figure 1. Constant current charge profiles of the LixNi0.8Co0.15Al0.05O2 (x = 0.5 and 0.1).
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RESULTS AND DISCUSSION Figure 1 shows the charge profile of a LixNi0.80Co0.15Al0.05O2 halfcell at a rate of C/10. We obtained half-charged (x = 0.5) and overcharged (x = 0.1) samples for subsequent examination. The crystal structure of pristine (before charging) LiNi0.80Co0.15Al0.05O2 powder was investigated as a reference. A number of particles were selected for observation and the results were consistent; thus, only a representative particle is shown in Figure 2. Figure 2a presents a HREM image of surface region, and the inset presents a BF image of the whole particle. A SADP [Figure 2b] obtained from the whole particle is indexed as a 11̅ 01 zone-axis of the layered structure (R3̅m). A fast Fourier
Figure 3. (a) High-resolution image of a half-charged NCA surface region. Inset displays the whole particle. (b) SADP from entire particle (i, ii) fast Fourier transformation results of area i and ii in (a), respectively. L and S indicate the layered and spinel structures, respectively.
Li0.5Ni0.80Co0.15Al0.05O2 particle. A BF image of the whole particle is included as an inset. A SADP from the whole particle [Figure 3b] indicates that it is oriented along the 1̅101 zone-axis of the layered R3̅m structure. In contrast, the local crystal structure as determined from the FFT [Figure 3i,ii] demonstrates the existence of other phases: FFTs from the near surface region (region i) are indexed as the 112 zone-axis orientation of the spinel structure. A pair of spots from Figure 3ii indicates an interplanar spacing of 0.23 nm corresponding to either the{11̅02} planes of the layered structure, the {222} planes of the spinel structure, or the {111} planes of the rock-salt structure. When the existence of the spinel structure at the near surface region is considered, it is unlikely that a layered structure is maintained at the outermost surface. In addition, the lattice parameters of the spinel and rock-salt structures are correlated with those of a layered structure in the hexagonal unit cell by the following relationship:12 c aspinel = ahex. × 2 2 = hex. (Eqn. 1) 3 arock salt = ahex. × Figure 2. (a) High-resolution image of the surface region of a pristine LiNi0.8Co0.15Al0.05O2 particle. The inset shows the overall sample. (b) Diffraction pattern of whole particle, indexed as the 1̅101 zone-axis orientation of layered R3̅m structure (L). (c) Fast Fourier transform from the region indicated in (a).
2 =
chex. 2 3
(Eqn. 2)
The lattice parameter of the spinel is twice as large as that of rocksalt structure; thus, the 111 reflection of the spinel structure [arrows in Figure 3i] is an indicator that allows us to determine whether the crystal structure is the spinel or the rock-salt. 1086
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We, thus, conclude that the spots in Figure 3ii originate from the rock-salt structure. According to these SADP and FFT results, three phases coexist in one particle after being brought to the half-charged state; the crystal structure of the bulk particle maintains the initial layered R3m ̅ structure, but phase transformations happen in the vicinity of surface, changing the layered structure to the disordered spinel immediately subsurface. Furthermore, another phase transformation at the surface results in the formation of the rock-salt structure within the first couple of nanometers of the surface. Prior work has suggested that severe electrochemical charging is needed to alter the surface structure of NCA.6,14 Thus, it is surprising that even a mild charge at room temperature can induce structural instability at the surface region. Wang et al. used first principle calculations to demonstrate that the phase transitions from the layered phase to the spinel or rocksalt phases are a kinetically controlled process and that the thermodynamic driving force for phase transitions is sufficiently large to favor these changes, even at 0 K.15 The three possible phases of NCA share the same oxygen framework, indicating that the phase transitions result from metal ion movements within an identical oxygen arrangement. Therefore, the temperature at which the phase transition occurs is related to the mobility of the metal ions and, thus, the amount of vacancies generated during charging.16 It is logical that the kinetics of lithium diffusion causes the extent of delithiation to vary as a function of the distance from the interface between the electrode and the electrolyte. Therefore, the phase transitions might happen more easily at the surface because there is a lower lithium content or higher number of lithium vacancies at the exterior of particle. This provides sufficient space within the lattice for the movement of the transition metal ions, even at room temperature. We also find that delithiation introduces significant structural variation in the overcharged NCA (Li0.1Ni0.80Co0.15Al0.05O2) particles. Because 90% of lithium ions are removed from the initial state, overcharged particles are very unstable and severe structural nonhomogeneity is found within a number of the particles. Figure 4a shows a BF image of an overall particle and an HREM image from the surface region. A SADP acquired from the whole particle [Figure 4b] is indexed as the layered R3̅m structure along the 1̅101 zone-axis; however, FFT results [Figure 4i,ii,iii] from the regions indicated in the HREM image [Figure 4a] demonstrate the existence of the rock-salt structure at the surface, with the subsurface structure being indexed as the spinel. Regardless of the state of charge, inhomogeneity in the crystal structure can be found within the vicinity of the surface of the particles, indicating that kinetic effects induced by lithium removal play an important role in the phase transitions. Overcharging also leads to a change in morphology. Figure 5 presents BF and HAADF images from an overcharged NCA particle. Both BF and STEM-HAADF images indicate that the outer area of particle has become porous. The exact thickness of the porous layer is hard to state in a systematic fashion as it is seen to vary substantially from particle to particle; however, it is clear that when 90% of the lithium is removed from the original cathode material, both the crystal structure and morphology are substantially modified at the surface area. The crystal structure for the layered R3̅m, spinel, and rock-salt structures each has the same close-packed oxygen lattice: that is, phase transformations take place by cation relocation within the same oxygen framework. These ionic movements result in a modification of the nearby coordination and bonding with oxygen, which is reflected in the electronic structure. XAS has been used to
Figure 4. (a) High-resolution image of an overcharged NCA surface region. The overall particle is shown as an inset. (b) SADP from whole particle, and (i−iii) fast Fourier transformation results of three regions in (a), respectively. L, S, and R denote the layered R3m ̅ , spinel, and rocksalt structures, respectively.
Figure 5. (a, b) BF images from the whole and surface area of overcharged NCA particle, respectively. (c, d) HAADF images from the whole and surface area of overcharged NCA particle, respectively.
determine the electronic structure of these cathode materials.13,14 Similar to XAS, EELS provides information about unoccupied density of states, but EELS has superior spatial resolution,17 1087
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Figure 6. BF images, SADPs, and EEL spectra of the oxygen K-edge from (a) pristine, (b) half-charged, and (c) overcharged NCA, respectively. Areas being analyzed are indicated with circles. Area i indicates the whole particle and areas ii and iii mean near-surface and edge area, respectively. The circles indicate the area from which the respective diffraction patterns and spectra were obtained.
display two distinct peaks, including a pre-edge peak at around 527−530 eV and primary peak at around 538.5−540 eV. The oxygen pre-edge peak is attributed to the transition of electrons from the 1s core state to unoccupied 2p states hybridized with 3d states in transition metals (in this case, mostly Ni).19 The main peak originates from a transition of electrons from the 1s state to hybridized states of O 2p and metal 4sp states.13 All spectra shown in Figure 6 are normalized to the intensity of the main peak. The SADPs in Figure 6a demonstrate that both the bulk and the surface of pristine NCA particles have the original layered R3m ̅ structure. In addition, the oxygen electronic structure is identical from the bulk to the surface. Thus, the crystallographic and electronic structures are uniform throughout particles before the electrochemical reaction.
allowing one to choose the nanoscale area from which the information is obtained.18 Furthermore, one can acquire an image, diffraction information, and an EEL spectrum from the same area, at nearly the same time. This ability to perform a simultaneous analysis with TEM/EELS enables us to understand comprehensively both the crystallographic and electronic structure modifications that occur within a nanoscale region. Figure 6 presents BF images of several particles, with indications of the areas being analyzed, representative oxygen K-edge EEL spectra, and diffraction patterns from pristine, half-charged, and overcharged NCA particles, respectively. Both spectra and diffraction patterns are obtained from the same area. Typical absorption spectra from the oxygen K-edge (Figure 6) of NCA 1088
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relative intensity of peaks in the EEL spectra can be modified by experimental effects such as sample thickness, which can lead to variation in the intensity of the pre-edge peaks. However, the spinel to rock-salt transition leads to substantial modifications to the chemical bonding between the oxygen ions and between the oxygen and transition metals ions,14 and thus should be reflected in the EEL spectra. Indeed, as lithium is removed from the original state, we find that oxygen pre-edge peak changes in important and interesting ways. Specifically, the intensity of the pre-edge peak increases at first but decreases thereafter with an accompanying shift to higher energy loss. Yoon et al. reported changes in the oxygen pre-edge peak of NCA cathode materials with charging through the use of X-ray absorption. Supporting Information Figure S1 is the reproduced oxygen K-edge XAS spectra using surface sensitive partial electron yield (PEY) mode from the reference.13 Data points were extracted from this reference through the use of a graphical digitizer application (GraphClick) and replotted in Supporting Information Figure S1 for clarity. Due to its superior energy resolution, XAS can discern three distinct oxygen preedge peaks: a peak at 528.5 eV (labeled as A) attributed to a transition of electrons from the 1s state to the hybridized O 2p−Ni3+/4+ 3d orbitals (consistent with the layered structure), a peak at 531.8 eV (B) originating from a transition of electrons from the 1s state to the hybridized state of O 2p−Ni2+ 3d orbitals (a NiO-like cubic phase), and a peak at 533.8 eV (C) resulting from the existence of an Li2CO3 phase on the surface.13,21 In comparison with the EEL spectra, XAS spectra have better energy resolution but no site specificity. Despite the poorer energy resolution of the EELS data, the trends in the pre-edges peaks of EELS and XAS are generally consistent. The increase in the pre-edge intensity can be interpreted as the result of charge compensation that occurs as lithium is extracted from the cathode material. Holes associated with the oxygen ions in O 2p−Ni 3d hybridized states take part in charge compensation because of the lithium deintercalation.22 Thus, as lithium is extracted the O 2p−Ni 3d hybridized states are increasingly empty and available for energy absorption, leading to an increase in the intensity of the pre-edge peak. With strong electrochemical reaction (i.e., overcharging), the pre-edge peak shifts to higher energy. The previous XAS data (Supporting Information Figure S1) demonstrates that the contribution of charge compensation is decreased by the O 2p−Ni3+/4+ 3d states but increased by hybridized O 2p−Ni2+ 3d states in the fully charged state (x = 0). Supporting Information Figure S2 presents O K-edge, Co L-edge, and Ni L-edge EEL spectra before and after electrochemical charging. After overcharging, distinct changes are noticed in a spectrum acquired with the smaller selective area aperture, which was chosen to select the surface area for analysis. Consistent with the oxygen K-edge modification, the Co and Ni L-edges have shifted to the lower energy loss, indicating that both transition metals have been reduced. Thus, our EELS results suggest that the transition metal ions at the surface that are bonded to oxygen are reduced when in the overcharged state (x = 0.1), resulting in a pre-edge peak shift to higher energy. Additionally, with the shift, the intensity of the pre-edge decreases considerably. In other words, the probability of an electron transition from the O 1s state to O 2p−Ni 3d hybridized state is lower, indicating that the effective electron density of oxygen increases concomitant with the increased reduction of the Ni ions. The extraction of lithium ions, the reduction of Ni ions, and the presence of a high effective electron density of oxygen ions are all in conflict with the thermodynamic need to
Figure 6b presents oxygen K-edge EEL spectra and corresponding SADPs from a half-charged (x = 0.5 in LixNi0.8Co0.15Al0.05O2) particle, specifically from (i) the entire particle, (ii) subsurface, and (iii) at the surface, respectively. The data from the entire particle [Figure 6b-i] tells us of the average structure. It should be noted that the relative intensity of the preedge peak to the main peak in EEL spectrum of the half-charged particle [Figure 6b-i] is slightly higher than that in EEL spectrum of the pristine particle [Figure 6a-i] due to the increased transition metal oxidation state and a participation of oxygen ions in charge compensation13 induced by electrochemical charging even though the bulk crystal structure is preserved as the layered structures. EELS and SADPs from the exterior region [Figure 6b-ii] are similar throughout the entire sample, indicating that the crystal and electronic structures are primarily in their initial, pristine state. However, both the EELS and SADP display abrupt changes at the outermost edge of the particle. At the outermost edge of the half-charged particle [Figure 6b-iii], the pre-edge peak shows a shift to higher energy, whereas the main peak shifts to lower energy, thereby making the change in the energy loss between the highest point of the pre-edge and main peak smaller. In addition, the intensity of pre-edge peak decreases. The SADP from the identical region indicates that the outermost surface of the particle has the rock-salt structure, as indicated by the 112̅ zone-axis pattern. There is, thus, a clear correlation between the phase transition to rock-salt and the change in the local electronic structure of oxygen. This is explored in greater detail below. In the case of overcharged NCA, the SADP from the whole particle shows diffraction spots from both the layered and spinel structures [Figure 6c-i]. Solely from the location of the diffraction spots shown in Figure 6c-i, the pattern can be interpreted as that from the disordered spinel structure, but comparing the intensity of each spot leads to the conclusion that there is a mixed structure consisting of both the layered and spinel structures. This indicates that a considerable portion of particle has undergone the first phase transition, with the outermost region changing from the layered structure to the spinel structure while maintaining a core of the original layered structure. It is interesting to note that even though a significant portion of the particle has undergone the first phase transition, the fingerprint of the EEL spectrum from the entire overcharged particle is comparable to that which we obtained from the bulk of the pristine sample, which did not undergo the phase transition. Figure 6c-ii displays SADP and EELS from the surface region of overcharged NCA. Crystallographic and electronic structural changes are prominent: the crystal structure has transformed into the rock-salt structure, the pre-edge peak of the EEL spectra has shifted to higher energy with a decreased intensity, and the main peak has shifted to a lower energy. The kinetics of the electrochemical reaction induces a gradient in the degree of lithium removal; this results in more substantial modification of both the crystallographic and electronic structures on the immediate surface area. It is notable that the electrochemical reaction induces a nascent phase transition and corresponding modifications in the oxygen K-edge EELS at the outermost surface of the half-charged NCA particle. With further charging, the crystallographic and electronic structure transformations extend further into the interior of the particles, leading to a substantial portion of the particles being modified. Furthermore, based on the relatively minor differences in the transition metal coordination around the oxygen atoms for the layered R3̅m and the disordered spinel,20 we would not anticipate significant changes in the oxygen K-edge. Notably, the 1089
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Figure 7. BF images and SADPs from the indicated regions of half-charged Li0.5Ni0.8Co0.15Al0.05O2 particles. L, S, and R indicate the layered, spinel, and rock-salt structures, respectively.
removed from the surface of a particle during charging, the greater number of lithium vacancies facilitates the phase transitions. Electrochemical delithiation of NCA is known to be a kinetically limited reaction; therefore, not only the microscopic (internal to the particles) but also the macroscopic degree of lithium removal could be different from particle to particle: the particle size effect shown in Figure 7 may be an example of a macroscopic kinetic effect. In order to more thoroughly investigate the particle-to-particle variation as a function of charge state, we acquired oxygen K-edge EEL spectra from ten different particles at each charging condition. We utilized two different sizes of selected area aperture to define the area from which the EEL spectra were obtained: a selected aperture of 325 nm of effective size within the image plane was used to obtain a bulk analysis of the particles, whereas the smallest available selective aperture, with a size of 180 nm at the image plane, was introduced to select only the surface of particles. Figure 8 summarizes these findings. The y axis plots the ratio of pre-edge peak intensity to that of the main peak (R) versus the magnitude of the energy loss difference between the highest point of the main peak and that of the pre-edge peak (ΔE). The magnitudes of R and ΔE values correlate to the degree of charge compensation by O 2p−Ni 3d hybridized states and the valence of Ni, respectively. Data from the bulk areas shows relatively small changes in the dispersion of either R or ΔE even after overcharging, with the exception of one peculiarly low data point with a ΔE of 7 eV. On the other hand, analysis with the smaller aperture shows that distribution of R or ΔE becomes significantly broader as charging proceeds, indicating that the surface of most particles sees significant modification after overcharging. The results from these oxygen K-edge EELS analysis are well matched with representative samples in Figure 6 and from the X-ray absorption result shown in Supporting Information Figure S1. However, it is noteworthy that the crystallographic and electronic structure modifications are not occurring in a completely uniform fashion. Considerable data from the bulk particles, as well as a few of the data points from the surface of the particles have comparable R and ΔE values to the pristine samples, indicating that the original electronic structure can be
maintain charge neutrality. This imbalance in charge results in the introduction of oxygen vacancies, leading to a loss of oxygen from the cathode material and the development of surface porosity, which is shown at Figure 5. The kinetics of lithium diffusion during the electrochemical reaction leads to a situation where the surface has lost more lithium than the bulk, provoking structural instabilities. Within the half-charged sample, incipient transformations are found at the very edge of particle [Figure 6b]. However, the degree of inhomogeneity is also different depending on the size of particles. Figure 7 presents BF images and SADPs from a half-charged NCA particle (Li0.5Ni0.80Co0.15Al0.05O2) at room temperature. Figure 7a indicates that interior of particle (region i) keeps its original layered structure while the first phase transition, from the layered to the spinel, occurs at the surface regions (ii, iii). A SADP from the inner region (i) in Figure 7a is indexed primarily as the layered R3̅m along the 1̅101 zone-axis. Additional faint spots in Figure 7a-i are attributed to the spinel structure, most likely originating from the near surface region. SADPs from regions ii and iii in Figure 7a are indexed as the 112̅ zone-axis orientation of the spinel structure, but spots forbidden by the structure factor are clearly seen (for example, the reflection from {420} planes). These 420 reflections (dotted circles) indicate that the spinel structure is in a disordered state. This disordered state arises because delithiation leads to the migration of transition metal ions, which in turn leads to the formation of an imperfect or disordered spinel structure.7 This disorder manifests itself as incomplete cancelations in the structure factor, leading to the presence of additional spots in the diffraction pattern. The influence of strong kinetic effects on the phase transformations is more evident for the smaller particles. Figure 7b shows a BF image and SADPs from a half-charged NCA particle with size of 500 nm, in comparison with the 1 μm-sized particle in Figure 7a. In this smaller particle, the phase transformation has egressed considerably into the interior of the particle, and the SADP of region Figure 7b-i indicates that there is a mix of phases: the layered R3̅m structure along the 2̅021 beam direction and the disordered spinel along the 011 zone-axis. It is interesting that a SADP from region ii indicates the formation of the rocksalt structure, with the 011 beam direction. In summary, our observations clearly indicate that as more lithium ions are 1090
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Figure 8. Statistical analysis of the relative strength of the pre-edge peak of the oxygen K-edge EEL spectra as a function of the charge state of the NCA particles.
Figure 9. Schematic summarizing the crystallographic and electronic structure changes that occur in LixNi0.8Co0.15Al0.05O2 cathode material upon electrochemical charging.
(the layered R3̅m structure), but phase transitions are observed at the exterior of the particle. Extreme extraction of lithium (90%) induces further phase transitions, making the core of the particle a mixed phase of the layered and spinel structures, with the outermost surface being transformed to the rock-salt phase. These structural changes are directly reflected in the EEL spectra. The EEL spectra show that the kinetics of delithiation lead to a higher probability of having a modified bonding configuration at the surface of the material. When lithium ions are extracted from the NCA cathode material, oxygen ions participate in the charge compensation, leading to readily discernible changes in the near-edge structure, in particular, when the phase transition to the rock-salt structure occurs. Depending on the degree of delithiation, the dominant mechanism is different; in the
maintained in the core of the sample as well as at specific locations on the surface. We summarize our conclusions with a schematic (Figure 9) that describes the crystallographic and electronic structure changes that occur in NCA during lithium deintercalation. We find that the local surface of the electrode is unstable against structural and chemical modifications even when the battery is charged at room temperature. The kinetics of lithium diffusion during charging induces an inhomogeneous state of charge inside of the particles: more lithium ions are removed at the surface, leading to phase transitions that take place as a result of the increased cation migration allowed by the excess vacancies introduced in the lattice. When 50% of the lithium is removed, the core of the particle maintains its original structure 1091
dx.doi.org/10.1021/cm403332s | Chem. Mater. 2014, 26, 1084−1092
Chemistry of Materials
Article
The authors also acknowledge support from the K-GRL Program funded by the Korea Institute of Science and Technology (Project No. 2Z04020). This work was supported by the National Research Foundation of Korea Grant funded by the Korean Government (MEST) (NRF-2011-C1AAA001-0030538).
half-charged state, the effective charge density is lower, whereas the overcharge state leads to a greater generation of oxygen vacancies. Furthermore, there is a correlation among the oxygen loss, the reduction of the transition metal ions, and the phase transition to the rock-salt structure. We find that both the degree of charging and the kinetics of charging govern the changes in crystallographic and electronic structure that occur in NCA. Because these structural instabilities are at the root of the detrimental figures of merit for battery systems, particularly for Ni-rich cathode materials, it is apparent that careful consideration of the state of the surface area is required to improve their performance. Our results suggest that some form of adequate surface coating may be one solution to improving both battery life and battery stability. Additionally, it is apparent that the strong kinetic effects we observe play a key role in determining the stability of the materials, indicating that an indepth analysis of the stability of these materials, especially the local surface area at the nanosale, as a function of temperature is merited, with the hope of finding configurations that promote kinetically limited reactions.
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CONCLUSION Analytical characterization with transmission electron microscopy has been used to delineate the inhomogeneity of crystallographic and electronic structure that occurs in the vicinity of LixNi0.8Co0.15Al0.05O2 particles surface as a function of the state of charge. Because of kinetic effects, there is a greater extraction of lithium from the surface of the particles, which leads to structural instabilities. These instabilities lead to the reduction of the transition metal ions, the loss of oxygen to maintain charge neutrality, and the formation of new phases as well as the porosity at the surface. These effects are more pronounced with an increasing degree of charge and are nonuniform due to both microscopic and macroscopic effects. Because of the observed, direct relationship between the degree of lithium extraction, the initiation of phase transitions, and the highly deleterious loss of oxygen from the structure, it is clear that improvements to the overall performance of these materials must occur through improvements in surface engineering as well is in enhanced control over the kinetics of lithium intercalation and deintercalation.
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ASSOCIATED CONTENT
S Supporting Information *
XAS of charged NCA, depending on the state of charge, taken from the literature and reproduced for clarity and exposition. O K-edge, Co L-edge, Ni L-edge EEL spectra from a pristine and the overcharged NCA. This material is available free of charge via the Internet at http://pubs.acs.org.
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AUTHOR INFORMATION
Corresponding Authors
*E-mail:
[email protected] (W. Chang). *E-mail:
[email protected] (E.A. Stach). Notes
The authors declare no competing financial interest.
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ACKNOWLEDGMENTS Research was carried out in large part at the Center for Functional Nanomaterials, Brookhaven National Laboratory, which is supported by the U.S. Department of Energy, Office of Basic Energy Sciences, under Contract No. DE-AC02-98CH10886. 1092
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