Correlating Structural Changes and Gas Evolution during the Thermal

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Correlating Structural Changes and Gas Evolution during the Thermal Decomposition of Charged LixNi0.8Co0.15Al0.05O2 Cathode Materials Seong-Min Bak,†,‡,∥ Kyung-Wan Nam,*,† Wonyoung Chang,‡ Xiqian Yu,† Enyuan Hu,† Sooyeon Hwang,§ Eric A. Stach,§ Kwang-Bum Kim,∥ Kyung Yoon Chung,*,‡ and Xiao-Qing Yang*,† †

Chemistry Department, Brookhaven National Laboratory, Upton, New York, 11973, United States Center for Energy Convergence, Korea Institute of Science and Technology (KIST), Seoul, 136-791, Republic of Korea § Center for Functional Nanomaterials, Brookhaven National Laboratory, Upton, New York, 11973, United States ∥ Department of Material Science and Engineering, Yonsei University, 134 Shinchon-dong, Seodaemoon-gu, Seoul, 120-749, Republic of Korea

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‡

S Supporting Information *

ABSTRACT: In this work, we present results from the application of a new in situ technique that combines timeresolved synchrotron X-ray diffraction and mass spectroscopy. We exploit this approach to provide direct correlation between structural changes and the evolution of gas that occurs during the thermal decomposition of (over)charged cathode materials used in lithium-ion batteries. Results from charged LixNi0.8Co0.15Al0.05O2 cathode materials indicate that the evolution of both O2 and CO2 gases are strongly related to phase transitions that occur during thermal decomposition, specifically from the layered structure (space group R3̅m) to the disordered spinel structure (Fd3̅m), and finally to the rocksalt structure (Fm3̅m). The state of charge also significantly affects both the structural changes and the evolution of oxygen as the temperature increases: the more extensive the charge, the lower the temperature of the phase transitions and the larger the oxygen release. Ex situ X-ray absorption spectroscopy (XAS) and in situ transmission electron microscopy (TEM) are also utilized to investigate the local structural and valence state changes in Ni and Co ions, and to characterize microscopic morphology changes. The combination of these advanced tools provides a unique approach to study fundamental aspects of the dynamic physical and chemical changes that occur during thermal decomposition of charged cathode materials in a systematic way. KEYWORDS: energy storage, safety, abuse tolerance, X-ray absorption spectroscopy, transmission electron microscopy

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release of oxygen-containing species (e.g., O2−, O−, O22−, and O2), which are highly reactive and can accelerate severe thermal runaway by reacting with the flammable electrolyte.8−11 It has been reported that the combustion of electrolyte by the released O2 is highly exothermic: the heat of this combustion reaction is an order of magnitude larger than the heat absorbed by the charged cathode materials during the thermal decomposition that resulted in the oxygen release,12 making the overall reaction highly exothermic. The link between phase transitions, oxygen release, and catastrophic failure motivates our studies here, and we exploit a new operando approach to investigate phase transitions and oxygen evolution in charged cathode materials, as a function of their charge state.

INTRODUCTION Global climate change and environmental concerns are major driving forces behind the intensified research and development of electric vehicles (EVs). Lithium-ion batteries (LIBs) are considered to be the best candidates to power them. However, in order to achieve large-scale commercialization of LIBs in transportation applications, several challenges must be addressed in terms of low cost, high energy density, long life, and improved safety.1−5 Specifically, improvements in the safety characteristics of LIBs are one of the most critical barriers to be overcome. These safety concerns are associated with the occurrence of exothermic reactions in charged batteries that result in thermal runaway and catastrophic failure.6,7 When in a highly delithiated state (i.e., “overcharged”), electrode materials become unstable and may degrade through exothermic or endothermic phase transitions.8−11 The resulting decomposition of the overcharged cathode materials can result in the © 2013 American Chemical Society

Received: September 25, 2012 Revised: December 20, 2012 Published: January 3, 2013 337

dx.doi.org/10.1021/cm303096e | Chem. Mater. 2013, 25, 337−351

Chemistry of Materials

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charging profile is shown in the Supporting Information (see Figure S1). The charged cells were then transferred to an argon-filled glovebox for disassembly. The charged cathode was thoroughly washed using DMC solvent to eliminate residual salts and then gently abraded from the current collector (Al foil) to create powder samples for the TR-XRD/MS, XAS, and TEM measurements. Combined TR-XRD and Mass Spectroscopy. The experimental setup of the in situ TR-XRD coupled with MS (i.e., TR-XRD/MS) is illustrated in Figure 1. The TR-XRD data were collected at Beamline

Many spectroscopic techniques have been developed and applied to the study of the thermal stability of charged cathode materials. These include time-resolved and high-temperature Xray diffraction (XRD),7,13−17 thermogravimetric analysis (TGA)/differential scanning calorimetry (DSC),18−20 and accelerating rate calorimetry (ARC).21−23 These approaches monitor the structural changes and/or heat generation that result during thermal decomposition. Mass spectroscopy (MS)24−26 has also been used alone, or in combination with TGA, for quantifying oxygen release. Although the intensive research using these in situ and ex situ characterization techniques has greatly enhanced our understanding of the thermal behavior of electrode materials, none of them alone can provide direct correlation between structural changes and oxygen evolution behavior that occurs during heating. Here, we report insights gained by combining in situ timeresolved X-ray diffraction (TR-XRD) and MS. This combination allows simultaneous observation of the structural changes and gas species that are evolved during the thermal decomposition of charged cathode materials.27 Although combined TR-XRD/MS has been employed to study catalysis previously,28,29 it has not been utilized in solid-state chemistry, especially in the structural study of battery materials. We have chosen to study LixNi0.8Co0.15Al0.05O2, since it is one of the high-energy-density cathode materials used in lithium-ion batteries. However, because of its high nickel content, it suffers from thermal instabilities that lead to changes in its crystal structure from the layered phase (R3̅m) to the spinel phase (Fd3̅m) to the rock-salt phase (Fm3̅m) with increasing temperature.27 The O2 release accompanied by these structural changes has been studied in prior work,30,31 but direct correlation between the two has not been reported. In addition, we use X-ray absorption spectroscopy (XAS) and environmental transmission electron microscopy (ETEM) to further elucidate specific aspects of the structural changes that occur in these same LixNi0.8Co0.15Al0.05O2 samples. XAS provides information regarding local structural changes surrounding certain transition-metal elements (Ni, Co), as well as changes in their oxidation state. ETEM provides nanoscale observation of the morphological changes that occur in individual particles during thermal decomposition. The combination of these advanced techniques provides a unique opportunity to understand the fundamental inter-relationships that occur during the thermal decomposition of the charged layered cathode materials. The information gained from this study provides insights into the rational design of advanced cathode materials and chemistries with improved thermal stability and high energy density.

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Figure 1. Setup for the combined in situ time-resolved X-ray diffraction (TR-XRD) and mass spectroscopy (MS) experiments. X7B (λ = 0.3184 Å) of the National Synchrotron Light Source (NSLS), using an image plate detector in the transmission mode. Approximately 3.5−4.0 mg of the charged cathode powder was loaded into a glass capillary with an inner diameter of 0.7 mm and two open ends. One end of the capillary was connected to He carrier gas (marked as inlet in Figure 1) and the other end (marked as outlet) was connected to a residual gas analyzer/mass spectrometer (Model RGA200, Stanford Research Systems) with a flow meter to detect gas species released from the sample during heating. Quartz wool was placed on each side of the sample to prevent movement of sample due to the He carrier gas flow. TR-XRD patterns (∼4 min for each XRD scan) and MS signals were simultaneously collected in a continuous manner as the sample was heated from room temperature to 500 °C for 4 h (i.e., at a heating rate of ∼2.0 °C min−1). For easy comparison with results in the literature, all the 2θ angles in this paper have been converted to values corresponding to the more common laboratory Cu Kα radiation (λ = 1.54 Å). For estimation of the lattice parameters, the XRD patterns of as-charged samples were refined by Rietveld methods using the GSAS package with the EXPGUI interface.32,33 Background, scale factor, zero point, lattice parameters, atomic positions, and coefficients for the peak shape function were iteratively refined until the fit converged. The refinement results are shown in Figure S2 in the Supporting Information. X-ray Absorption Spectroscopy. Ni and Co K-edge XAS spectra were collected at the X18A and X19A beamlines of the NSLS using a Si (111) double-crystal monochromator in a transmission mode. The beam intensity was detuned to 70% of its initial intensity to minimize the high-order harmonics. Reference spectra of Ni and Co metallic foils were collected simultaneously with all of the spectra for energy calibrations. The ex situ XAS spectra were collected at room temperature after heating the charged cathode electrodes at each specific elevated temperature between 25 °C and 450 °C. The charged cathode electrodes for the ex situ XAS measurements were first washed by DMC solvent and then heated and cooled in a furnace under a He gas flow. The charged cathodes were sealed with Kapton tape immediately after the heat treatment, to prevent air exposure. The Ni and Co K-edge XAS spectra were processed using the Athena and Artemis programs.34 For each spectrum, the AUTOBK code was used to normalize the absorption coefficient, μ(k), and separate the χ(k) functions from the isolated atom absorption background. The

EXPERIMENTAL SECTION

Electrode Preparation and Electrochemical Tests. Charged cathode materials (LixNi0.8Co0.15Al0.05O2, from Ecopro Co. Ltd., Korea) with lithium compositions of x = 0.5, 0.33, and 0.1 were prepared by electrochemical delithiation using a constant current (i.e., galvanostatic) charge at a C/18 rate. The cathode electrodes were prepared from a mixed slurry of 80 wt % active material, 10 wt % carbon black (Chevron), and 10 wt % of PVDF (Kureha) as a binder in a n-methyl pyrrolidone (NMP) solvent. This slurry was subsequently coated onto an Al foil. The cathode materials were incorporated into 2032 coin cells with a Li metal foil anode, a Celgard separator, and an electrolyte of 1.2 M LiPF6 dissolved in ethylene carbonate (EC) and dimethyl carbonate (DMC) solvent (3:7 by volume). The lithium contents were estimated by the charge passed in the cell, assuming 100% columbic efficiency. The galvanostatic 338



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Figure 2. (a) TR-XRD patterns and simultaneously measured mass spectra (MS) for (b) O2 and (c) CO2, released from Li0.5Ni0.8Co0.15Al0.05O2 during heating to 500 °C. The left panel shows the models of ideal crystals with rhombohedral, spinel, and rock-salt structures.

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photoelectron energy origin (E0) is chosen at the first inflection point of the absorption edge jump. The extracted extended X-ray absorption fine structure (EXAFS) signal, χ(k), was weighted by k2 to emphasize the high-energy oscillations and then Fourier-transformed in k-ranges of 3.4−11.0 Å−1 for Co and 3.5−13.5 Å−1 for Ni, using a Hanning window function to obtain the magnitude plots of the EXAFS spectra in R-space (Å). The Fourier-transformed peaks were not phasecorrected; therefore, the actual bond lengths are ∼0.2−0.4 Å longer. The filtered Fourier transforms of EXAFS spectra in a R-range of 1.0− 2.9 Å were then fitted using theoretical single scattering paths generated with the FEFF 6.0 ab initio simulation code, using a rocksalt model structure. The amplitude reduction factors (S02) were determined to be 0.84 for Co and 0.75 for Ni from the preliminary fitting sessions, and then fixed during the final fitting, unless noted otherwise. The same inner shell potential shift (ΔE) was shared for the M−O and M−M shells (M = Co and Ni), while separate fitting parameters of the bond distance (r) and Debye−Waller factor (i.e., mean square disorder, σ2) were used for each shell. The coordination number for the first M−O shell was fixed to 6 (i.e., octahedral coordination) while those of the second M−M shell were refined during fitting. In Situ TEM. Nanoscale morphological changes occurring in the charged cathode material during heating were evaluated through the use of in situ TEM. The samples were investigated in real time during heating, over the temperature range of room temperature to ∼400 °C, using a Gatan 652 double-tilt heating holder and a spherical aberration corrected environmental TEM (FEI Titan 80/300 ETEM, with CEOS Cs-corrector) at an accelerating voltage of 300 kV. For the TEM experiments, the charged cathode particles were abraded from the current collector and dispersed in a DMC solvent in a small glass bottle. Supersonic vibration was applied to ensure well-dispersed particles before dropping the solution onto a standard TEM holey carbon grid. All TEM samples were prepared and loaded into the heating holder in a glovebox, and then transferred to the microscope in a hermetically sealed container to prevent sample exposure to air. For in situ heating TEM observations, the sample was brought to the target temperature and held constant for ∼30 min, allowing the sample stage to achieve thermal equilibrium prior to data acquisition.

RESULTS AND DISCUSSION

Combined Time-Resolved X-ray Diffraction (TR-XRD) and Mass Spectroscopy (MS). Figure 2 presents results from the combined TR-XRD/MS characterization of charged LixNi0.8Co0.15Al0.05O2 with x = 0.5 (cutoff voltage of 3.91 V vs Li/Li+; see Figure S1 in the Supporting Information). The evolution of both O2 and CO2 gases are observed to be closely related to the phase transitions that occur during thermal decomposition. Figure 2a displays a series of TR-XRD patterns of the Li0.5Ni0.8Co0.15Al0.05O2 cathode material recorded during heating from 25 °C to 500 °C (at a rate of ∼2 °C min−1), under a He carrier gas flow. The corresponding profiles for the oxygen (O2, m/z = 32) and carbon dioxide (CO2, m/z = 44) gas release obtained by MS are shown in Figures 2b and 2c. The XRD pattern for the as-charged Li0.5Ni0.8Co0.15Al0.05O2 indicates a rhombohedral (space group R3̅m) structure with the lattice parameters a = 2.8317(4) Å and c = 14.4085(1) Å, calculated by Rietveld refinement (see Figure S2 in the Supporting Information). The phase transition from the rhombohedral to disordered spinel (space group Fd3̅m) phase starts at ca. 194 °C, as indicated by the shift in position of the (108)R and (110)R peaks in opposite directions. The coalescence of these (108)R and (110)R peaks in the layered structure indicates the onset of the formation of the disordered spinel phase, and occurs through cation migration.30,31,35 The evolution of the (220)S diffraction peak at 2θ ≈ 31°, which is prohibited by the rhombohedral symmetry, also confirms the formation of the disordered spinel phase.30,31 This disordered spinel phase transformed gradually to the rock-salt phase as the temperature increased from 275 °C to 500 °C, as is evident from the decrease in the intensity of the spinel (311)S and (511)S peaks. At the end of heating at 500 °C, the main peaks were indexed to be a cubic rock-salt phase (space group Fm3̅m), with some minor peaks corresponding to the residual of the disordered spinel phase. 339



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Figure 3. (a) TR-XRD patterns and simultaneously measured mass spectra (MS) for (b) O2, and (c) CO2, released from Li0.33Ni0.8Co0.15Al0.05O2 during heating to 500 °C.

1 {Li1.0(M3.5 +)2 O4 } (spinel, Fd 3̅ m) 2 1 1 → {Li1.0(M 2.5 +)2.0 O3.0 } (rock salt, Fm 3̅ m) + O2 ↑ 2 4

During each phase transition, two O2 release peaks were observed, with the maximum evolution occurring at ∼240 and 320 °C in the MS profile, shown in Figure 2b. Generally, these O2 release peaks are clearly correlated to the two structural phase transitions, from the layered to the spinel and the spinel to the rock-salt phase. The beginning of the first noticeable O2 release observed at ca. 200 °C is very close to the temperature where the structural transition from the initial layered structure to the disordered spinel phase starts. After the first O2 release peak reached its maximum intensity at ca. 230 °C, a slight decrease of O2 occurred, followed by a second O2 peak that was initiated at ca. 275 °C and reached a maximum at ca. 320 °C. After that, a continuous O2 release was observed, but with a decreasing rate, until the end of the experiment at 500 °C. The onset temperature of the second O2 release at 275 °C is wellcorrelated to the second phase transition from the disordered spinel structure to the rock-salt phase. In addition, the temperature range of the second O2 release peak between 275 and 500 °C is also coincident with the temperature range of the second phase transition. A larger total amount of O2 gas was released during the second phase transition than the first one. Note that the O2 loss during the first phase transition from the layered to the disordered spinel phase was unexpected, based on the previous TGA/MS results on the charged cathode materials such as Li 0.5 Ni 1.02 O 2 , Li 0.5 Ni 0.89 Al 0.16 O 2 , and Li0.5Ni0.70Co0.15Al0.15O2.30,31 Such O2 release should not happen for the Li0.5Ni0.8Co0.15Al0.05O2 sample at this Li content, if it were to occur via the following homogeneous reaction (eq 1): Li1/2(M 3.5 +)1.0 O2 (layered, R 3̅ m) → (spinel, Fd 3̅ m),

no oxygen loss

(2)

where “layered” = LiMO2, “spinel” = LiM2O4, and “rock salt” = (Li +M)O. In the case of the LixNi0.8Co0.15Al0.05O2 with x = 0.5 (50% lithium extraction), the first phase transition, corresponding to the progressive transition between the layered and disordered spinel structure described in the reaction (eq 1), is believed to occur through cation migration (e.g., Ni, Co, and Al) from the original layer to the adjacent Li layer and through the displacement of the Li+ ions from octahedral to tetrahedral sites, without any oxygen loss from the lattice.30 This will be explored in further detail through the use of X-ray absorption spectroscopy (XAS) below. Based on eq 2, oxygen release is only anticipated to occur during the second phase transition over 275 °C: this agrees well with the results shown in Figure 2. However, the experimental observation of O2 release during the first phase transition of the Li0.5Ni0.8Co0.15Al0.05O2 strongly suggests that some parts of the sample (likely near the surface) are in a more delithiated state, compared with the average delithiation state of x = 0.5. We suspect that this inhomogeneity in the lithium composition of the charged sample stems from kinetic effects during charging. Although we used a relatively slow charge rate of C/18 for the delithiation, the surface of the particle may be more delithiated than the average lithium content of x = 0.5, because of the slow lithium diffusion kinetics. Thus, the more-delithiated parts could undergo a slight O2 loss reaction, even in the first phase transition region. This result implies that high rate charging and discharging of LIBs can considerably affect the safety characteristics of the cell by creating inhomogeneous regions near the surface. There is also

1 {Li1.0(M3.5 +)2.0 O4 } 2 (1) 340



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Figure 4. (a) TR-XRD patterns and simultaneously acquired mass spectra (MS) for (b) O2, and (c) CO2, released from Li0.1Ni0.8Co0.15Al0.05O2 during heating to 500 °C.

a possibility that the O2 release between 200 °C and 260 °C is due to the residual DMC solvent used for washing the samples, which was reported to decompose at ∼250 °C.36 In order to rule out this artifact, TR-XRD/MS data were collected for a pristine (i.e., uncharged) Li1.0Ni0.8Co0.15Al0.05O2 sample washed with the DMC solvent under the exact same condition applied to the washing of charged samples. The results showed no structural changes in the TR-XRD and no oxygen gas release up to 500 °C (see Figure S3 in the Supporting Information), suggesting that the amount of residual DMC is negligible after washing, if there is any, and does not affect O2 evolution in our case. This confirms that the oxygen release observed for the charged samples originates from the thermal decomposition of the cathode materials, not from the residual DMC. In addition to the O2 gas release, we were also able to monitor the release of carbon dioxide (CO2) during heating. The formation of CO2 is associated with the oxidation of carbon (from either the PVDF binder or the conducting carbon in the charged electrode) by the released oxygen. As shown in Figure 2c, the broad CO2 release peak started at ca. 250 °C and continued to 500 °C, with a maximum at ∼380 °C, which occurred at a higher temperature than the second O2 release peak at 275 °C in Figure 2b. It has been reported in the literature that the temperatures for CO2 release caused by the thermal decomposition of the PVDF binder alone (not in the electrode) under inert gas flow are in the range of 450−650 °C.37−39 It is also well-known that the conducting carbon materials are quite stable to temperatures as high as 700 °C under an inert gas flow.40−42 Therefore, the CO2 release at temperatures between 250 and 500 °C observed in Figure 2c is likely to be the result of an oxidation reaction of the carbon with the oxygen released from cathode materials, but not from O2 in the atmosphere. This emphasizes that the oxygen released from the charged cathode is very reactive and, thus, can seriously affect the safety of LIBs.

As shown in the TR-XRD patterns in Figure 3a, the LixNi0.8Co0.15Al0.05O2 with x = 0.33 (cutoff voltage at 4.12 V vs Li/Li+; see Figure S1 in the Supporting Information) goes through a similar set of phase transitions as observed for the x = 0.5 sample. However, the phase transitions start at lower temperatures and occur over a narrower temperature range for the disordered spinel phase, in comparison with that observed for the Li0.5Ni0.8Co0.15Al0.05O2 sample. The first phase transition (layered → disordered spinel) starts at ca. 186 °C and is completed at ca. 234 °C. This is followed by the second phase transition (disordered spinel → rock salt), which is completed at 476 °C. In contrast to the Li0.5Ni0.8Co0.15Al0.05O2 sample, which showed the existence of a residual spinel phase, the Li0.33Ni0.8Co0.15Al0.05O2 sample transformed completely to the rock-salt phase at the end of the thermal cycle. This suggests that, as the amount of charge is increased (i.e., when the material has a lower Li content), the materials become less thermally stable. The corresponding O2 release observed by mass spectroscopy is plotted in Figure 3b. Again, there are two O2 release peaks that are closely correlated with the two phase transitions observed in the TR-XRD patterns in Figure 3a. In comparison with the O 2 release behavior for the Li0.5Ni0.8Co0.15Al0.05O2 sample, there is a lower onset temperature (i.e., 180 °C) for the first O2 release peak, which agrees well with the TR-XRD results in Figure 3a. The temperature range of the second O2 peak is also well-correlated to the second phase transition region between 250 °C and 500 °C. The relatively larger amount of O2 release for the second reaction, in comparison with the Li0.33Ni0.8Co0.15Al0.05O2 sample, is described well by the following reactions: Li1/3(M 3.67 +)1.0 O2 (layered, R 3̅ m) 4 1 → {Li3/4(M 3.21 +)9/4 O4 } (spinel, Fd 3̅ m) + O2 ↑ 9 9 (3) 341



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Figure 5. Contour plot of the intensity of the (220)s XRD peak versus 2θ (28°−34°) and temperature (25−500 °C) for (a) Li0.5Ni0.8Co0.15Al0.05O2, (b) Li0.33Ni0.8Co0.15Al0.05O2, and (c) Li0.1Ni0.8Co0.15Al0.05O2.

4 {Li3/4(M 3.21 +)9/4 O4 } (spinel, Fd 3̅ m) 9 4 2 → {Li3/4(M 2.33 +)9/4 O3} (rock salt, Fm 3̅ m) + O2 ↑ 9 9

layered to the disordered spinel. After the rapid O2 release, the second broad O2 release peak occurred over a wide temperature range, from 242 °C to 405 °C. In this sample, more O2 gas was released during the first phase transition than the second one, which is consistent with the following theoretical calculations:

(4)

where “layered” = LiMO2, “spinel” = (Li + M)3O4, and “rock salt” = (Li +M)O. Figure 3c shows the CO2 release from the Li0.33Ni0.8Co0.15Al0.05O2 sample during thermal decomposition. It should be noted that a small CO2 release peak is clearly observed in the lower temperature region between 200 and 250 °C, corresponding to the first phase transition from the layered phase to the disordered spinel phase. This is in addition to the main CO2 release peak, which appeared at ∼250 °C and continued up to 500 °C. Again, this is well-correlated to the temperature range of the second phase transition. The TR-XRD patterns and corresponding MS profile of the O 2 and CO 2 gases for t he h ig hly overch arged LixNi0.8Co0.15Al0.05O2 sample with x = 0.1 (cutoff voltage at 4.77 V vs Li/Li+; see Figure S1 in the Supporting Information) are plotted in Figures 4a, 4b, and 4c. In comparison with the Li0.5Ni0.8Co0.15Al0.05O2 and Li0.33Ni0.8Co0.15Al0.05O2 samples, the phase transitions initiated at lower temperatures for both the first and second phase transitions, and a much narrower temperature range was observed for the disordered spinel phase (Figure 4a). The first phase transition starts at ca. 178 °C and is completed at ca. 215 °C, while the second phase transition starts at 215 °C and continues to ∼400 °C. Several interesting features were observed in this highly overcharged Li0.1Ni0.8Co0.15Al0.05O2 sample. The rapid structural changes resulted in a very sharp O2 release peak, with a maximum at ∼205 °C, as shown in Figure 4b. This release starts at 170 °C, which is close to the temperature of structural change from the

Li1/10(M 3.9 +)1.0 O2 (layered, R 3̅ m) 11 {Li3/11(M 2.84 +)30/11O4.0 } (spinel, Fd 3̅ m) → 30 4 O2 ↑ + 15

(5)

11 {Li3/11(M2.84 +)30/11O4.0 } (spinel, Fd 3̅ m) 30 11 {Li3/11(M 2.84 +)30/11O3.0 } (rock salt, Fm 3̅ m) → 30 11 O2 ↑ + 60

(6)

where “layered” = LiMO2, “spinel” = (Li + M)3O4, and “rock salt” = (Li +M)O. Interestingly, a significant, sharp CO2 release peak was also observed, which initiated at ca. 175 °C, and had a maximum at ∼197 °C, as shown Figure 4c. The fact that CO2 was evolved at such a low temperature is a clear evidence of the highly active nature of the oxygen species (such as O2−, O−, O22−) released from the overcharged Li0.1Ni0.8Co0.15Al0.05O2. These highly active oxygen species are responsible for accelerating the oxidation of the binder and/or the conducting carbon in the electrode and thus decrease the temperature of the associated CO2 release.43,44 To compare the O2 release behavior of each sample, the MS results for O2 release are plotted together in Figure S4 in the Supporting Information. The more significantly the materials are overcharged, the lower the O2 342



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Figure 6. Ni K-edge XANES spectra of (a) Li0.5Ni0.8Co0.15Al0.05O2, (b) Li0.33Ni0.8Co0.15Al0.05O2, and, (c) Li0.1Ni0.8Co0.15Al0.05O2 after heating at various temperatures. Insets show the detailed features of their pre-edge region. (d) Variation of the Ni K-edge positions at the first inflection point for both cathodes as a function of temperature. The edge position for the reference NiO (Ni2+) spectrum is also marked as a broken line.

temperature range are critical to improving the safety characteristics of LIBs. The phase transformations observed during the thermal decomposition are known to be closely associated with the migration of transition-metal cations and, in turn, to the release of oxygen during heating. Therefore, it is important to understand this cation migration mechanism in greater detail. Figure 5 presents the diffraction intensity contours for the (220)S peak at 2θ ≈ 31° in the TR-XRD patterns for the

release temperature and sharper the O2 release peak. In particular, the sharp O2 release at low temperature would definitely cause serious safety problems. In a real Li+-ion cell, where a large amount of highly reductive electrolyte is available, the pulse of highly active oxygen species released from the cathode might quickly react with the flammable electrolyte and accelerate the thermal runaway. Our results indicate that suppressing the oxygen release, pushing its starting point to higher temperatures, and spreading its reaction to a wider 343



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Figure 7. Co K-edge XANES spectra of (a) Li0.5Ni0.8Co0.15Al0.05O2, (b) Li0.33Ni0.8Co0.15Al0.05O2, and (c) Li0.1Ni0.8Co0.15Al0.05O2 after heating at various temperatures. The insets in these figures show detailed features of the pre-edge region. (d) Variation of the Co K-edge positions at the first inflection point for both cathodes as a function of temperature. The edge positions for the reference CoO (Co2+) and Co3O4 (Co2.67+) spectra are also marked as broken lines. Panels (e), (f), and (g) show contour intensity maps of the pre-edge spectra with increasing temperature in the insets of Figures 7a, 7b, and 7c), respectively.

∼250 °C and reaches a maximum at ca. 400 °C; it then slightly decreases during further heating. This suggests that the TMs first migrate to the tetrahedral sites to form the M3O4-type spinel structure and then move to the adjacent octahedral sites to form the rock-salt-type structure. As shown in Figures 5b and 5c, the maximum intensity of the (220)S peak was observed at ∼300 and 225 °C for the Li0.33Ni0.8Co0.15Al0.05O2 and Li0.1Ni0.8Co0.15Al0.05O2 samples, respectively, which is much lower than the Li0.5Ni0.8Co0.15Al0.05O2 case. Thus, it is clear that, for the more highly charged samples, the (220)S peak exists over a narrower temperature region and has a maximum intensity at a lower temperature. This indicates a poorer thermal stability, with faster cation migration between two octahedral sites through the intermediate tetrahedral sites. This

LixNi0.8Co0.15Al0.05O2 (x = 0.5, 0.33, and 0.1) with heating temperature. Because the (220)S spinel peak is sensitive to partial occupancy of the transition metals (e.g., Ni and Co) in the 8a tetrahedral sites, the intensity of this peak should be a reliable indicator of cation migration into the tetrahedral sites. If a significant amount of transition metal (TM) cations occupy the tetrahedral sites, the structure is the M3O4-type spinel. This structure differs from the LiMn2O4-type spinel structure, where most of the tetrahedral sites are occupied by Li. It is known that the higher and wider temperature range of the M3O4 (e.g., Co3O4)-type spinel structure pushes the phase transition to the rock-salt structure to a higher temperature.27 For the Li0.5Ni0.8Co0.15Al0.05O2 sample, as can be seen from Figure 5a, the intensity of the (220)s peak starts to increase at 344



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sample within a similar narrow temperature range (shown in the MS result of Figure 4b). Co K-edge XANES spectra in Figure 7 also show similar shifts in edge positions toward lower energies with temperature, as were observed in the Ni K-edge, indicating that thermal decomposition occurs around Co as well. The Co spectra show that the entire edge shifts with heat treatment up to 450 °C, as seen in the main peak region (labeled “B”). In addition, a more significant reduction occurs in the case of the Li0.1Ni0.8Co0.15Al0.05O2 sample than in the less-charged ones, as seen in Figure 7d. This plot presents variations of the edge positions at the first inflection point in the Co K-edge XANES spectra, as a function of temperature. However, when compared with the Ni K-edge shifts in Figure 6d, Co edges begin to shift at slightly higher temperatures and to a lesser extent than the Ni edges for all of the charged LixNi0.8Co0.15Al0.05O2 samples. For example, Co ions in the less-charged Li0.5Ni0.8Co0.15Al0.05O2 and Li0.33Ni0.8Co0.15Al0.05O2 samples were not fully reduced to Co2+ and maintained their oxidation state close to Co2.67+ (i.e., Co3O4) at 450 °C, whereas the Ni ions fully reduced to Ni2+ (i.e., NiO) at the same states of charge (Figure 6d). This indicates that the Co reduction and migration follow the first Ni reduction and migration during thermal decomposition. This indicates that Co ions are more thermally stable than Ni ions. Interestingly, both the Co pre-edge peak intensity (shown in the insets of Figures 7a−7c) and corresponding contour intensity maps (Figures 7e−7g) clearly show nonmonotonic changes with increasing temperature. For all of the LixNi0.8Co0.15Al0.05O2 samples, the Co pre-edge peak intensity first increased with increasing temperature, reached a maximum intensity, and then decreased during further heating to 450 °C. The maximum intensity of the Co pre-edge peak for the Li 0.5 Ni 0. 8 Co 0 .15 Al 0.0 5 O 2 , Li 0.3 3 Ni 0 .8 Co 0.15 Al 0. 05 O 2 , and Li0.1Ni0.8Co0.15Al0.05O2 samples occur at ca. 350, 300, and 250 °C, respectively. This is quite different than that observed in the Ni pre-edge peak features in the same samples. For nickel, there was a continuous decrease in the intensity (insets of Figures 6a−c), because of the continuous reduction of Ni ions, without a change in their coordination environment (i.e., they maintained octahedral coordination). In contrast, the observation of an increasing intensity of the Co pre-edge in the early stages of heating indicates that the some Co ions in LixNi0.8Co0.15Al0.05O2 sample migrate into the adjacent tetrahedral sites first and tend to form a Co3O4 spinel-like local structure (1/3 of Co ions reside in the tetrahedral sites in the Co3O4 spinel). Note that this structure is considered to be highly disordered and nonstoichiometric with defects and different cation distributions, in comparison with stoichiometric Co3O4 spinel. This is due to the loss of oxygen and the existence of Li, Ni, and Al ions. A decrease in the average oxidation state of Co ions close to 2.67+ (Co3O4) for the charged samples is seen in Figure 7d. This also supports the conclusion that a Co3O4 spinel-like structure is formed around Co. After the maximum peak intensity is reached, the decreasing intensity of Co pre-edge indicates that there is further migration of Co ions from the tetrahedra to adjacent octahedral sites (which were originally occupied by Li before charging); this occurs during the phase transition from the disordered spinel to the rock-salt phase. For the more highly charged samples of LixNi0.8Co0.15Al0.05O2 (with x = 0.33 and 0.1), the pre-edge peak maxima were reached at lower temperatures and within a narrower temperature range. This indicates a more rapid phase transition from the disordered

result also agrees well with the lower oxygen release temperatures which occur over a narrower temperature region for the more highly charged samples (shown in Figure S4 in the Supporting Information). This type of cation migration strongly depends on the oxidation state of the cations and their environment, as well as the oxygen skeleton of the structure.45 We explore this further using XAS below, since this technique is capable of probing the local structural changes surrounding certain TMs (e.g., Ni and Co), as well as changes in their oxidation state. X-ray Absorption Near-Edge Structure (XANES). Ni Kedge X-ray absorption near-edge structure (XANES) spectra for the charged LixNi0.8Co0.15Al0.05O2 cathodes with x = 0.5, 0.33, and 0.1 after heat treatment up to 450 °C are shown in Figures 6a, 6b, and 6c. In all spectra, two distinct edge features are observed and are marked as “A” and “B”. The weak absorption pre-edge peaks marked as “A” (which are also shown in detail in the inset of each figure) are associated with the dipoleforbidden 1s→3d electronic transition. However, this transition becomes partially allowed when hybridization of the 3d and 4p orbital occurs, because of structural distortions in the local symmetry between the TM and oxygen coordination.46 Therefore, TMs in tetrahedral sites usually show much stronger pre-edge peak intensity than in the octahedral sites, because of the decrease of an inversion center in the tetrahedral coordination symmetry.47 The strong main absorption peak, marked as “B”, originates from the electronic dipole-allowed transition of a 1s core electron to an unoccupied 4p orbital.46 Ni K-edge XANES spectra indicate a clear edge shift toward lower energies in both the pre-edge and the main-edge onset in all samples with increasing temperature. This suggests that there is a reduction of the average oxidation state of Ni ions during heating, which, in turn, contributes to the thermal decomposition in the charged cathodes. During heating, the edge shift initiates above 200 °C (dark green line) in the case of the Li0.5Ni0.8Co0.15Al0.05O2 and Li0.33Ni0.8Co0.15Al0.05O2, but between 150 °C (cyan line) and 200 °C in the case of the more highly charged Li0.1Ni0.8Co0.15Al0.05O2. This is in good agreement with the lower onset temperature of the observed first phase transition for the highly charged Li0.1Ni0.8Co0.15Al0.05O2, in comparison with the two lesscharged materials. This indicates that Ni migration into the lithium layer is mostly responsible for the first phase transition. In addition, these edge shifts indicate that the reduction of Ni3+/4+ ions to Ni2+ in the charged LixNi0.8Co0.15Al0.05O2 occurs during heating; this would drive the oxygen release in order to satisfy charge neutrality. Also note that the intensity of Ni preedge remains quite low at all temperatures, suggesting that the Ni ions remain in octahedral sites (either in their original sites or in Li layer sites) during thermal decomposition. For a better comparison of the average oxidation state changes of Ni ions, variations of the edge positions at the first inflection point in the Ni K-edge XANES spectra are plotted in Figure 6d, as a function of heating temperature. The Li0.1Ni0.8Co0.15Al0.05O2 sample shows the largest edge shift (∼3.4 eV up to 450 °C) and the lowest onset temperature (below 200 °C) for Ni reduction (i.e., thermal decomposition of Ni ions), when compared with less-charged samples (x = 0.33, 0.5). In addition, ca. 70% of this reduction to Ni2+ is completed for the Li0.1Ni0.8Co0.15Al0.05O2 sample within a very narrow temperature range (between 150 and 250 °C). This agrees well with the massive oxygen release from the same 345



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Figure 8. (a) Crystal structure of (left) ordered rock-salt (e.g., layered) and (right) disordered rock-salt (e.g., NaCl-type rock-salt) model structures; (b) Fourier transformed magnitude of Ni K-edge EXAFS spectra for the charged Li0.5Ni0.8Co0.15Al0.05O2 (black) and Li0.1Ni0.8Co0.15Al0.05O2 (red) after heating at various temperatures in comparison with the reference NiO (bottom); (c) corresponding refined structural parameters; bond lengths for the first Ni−O shell (subpanel (i)) and second Ni−M shell (subpanel (ii)); Debye−Waller factors for the first Ni−O shell (subpanel (iii)) and second Ni-M shell (subpanel (iv)); and coordination number for the second Ni−M shell (subpanel (v)).

rock-salt structure with ordering of Li and M (M = Ni and Co) along alternating (111) planes. If there is no ordering of Li and M within a layer (i.e., complete cation mixing), the NaCl-type rock-salt structure is formed, as shown in Figure 8a (right). The coordination numbers of the first and second neighboring M− O and M−M shells in the stoichiometric rock-salt structures CoO and NiO are 6 and 12, respectively (Figure 8a, right). In contrast, in the ordered rock-salt structure (e.g., layered LixNi0.8Co0.15Al0.05O2) the same 6-fold coordination is observed for both the first and second M−O and M−M shells (Figure 8a, left). Therefore, the variation in coordination number for the second M−M shell is a good indicator of the degree of cation mixing and the formation of the NaCl-type rock-salt phase. In this regard, cubic rock-salt model structures (e.g., NiO and CoO) were used to generate theoretical paths for the first two M−O and M−M shells in order to fit the EXAFS of the Ni and Co K-edges. The coordination number for the first M−O shell was fixed to 6 (i.e., octahedral coordination) while that of the second M−M shell was refined during fitting. The bond length (r) and Debye−Waller factor (σ2) of each shell were refined

spinel to the rock-salt structure, resulting in a more-rapid oxygen release at lower temperatures. The fact that the changes in the Co pre-edge intensity (Figure 7) during heating are very similar to the (220)s peak intensity changes in the TR-XRD (Figure 5) also confirms that the initial migration of Co ions to tetrahedral sites is followed by migration to the octahedral sites. These results not only further support our conclusions regarding the pathways of the phase transitions observed in the TR-XRD results shown in Figures 5a−c, but also provide more detailed information about the different migration pathways of Ni and Co cations and their roles in driving the phase transitions and oxygen release. Extended X-ray Absorption Fine Structure (EXAFS). To identify the local structural variations around the Ni and Co sites during thermal decomposition, EXAFS spectra from the Ni and Co K-edges for the charged LixNi0.8Co0.15Al0.05O2 (x = 0.5 and 0.1) samples were collected and analyzed. The results are shown in Figures 8 and 9, as a function of heating temperature. As depicted in Figure 8a (left), the layered LixNi0.8Co0.15Al0.05O2 structure can also be considered as a 346



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Figure 9. (a) Fourier-transformed magnitude of Co K-edge EXAFS spectra for the charged Li0.5Ni0.8Co0.15Al0.05O2 (black) and Li0.1Ni0.8Co0.15Al0.05O2 (red) after heating, and in comparison with the reference CoO and Co3O4 (bottom); (b) corresponding refined structural parameters; bond lengths for the first Co−O shell (subpanel (i)) and second Co−M shells (subpanel (ii)); Debye−Waller factors for the first Co−O shells (subpanel (iii)) and second Co−M shells (subpanel (iv)); and coordination number for the second Co−M shell (subpanel (v)).

with the single scattering algorithm. Detailed fitting results and the EXAFS reliability R-factor are shown in Figures S5−S7 in the Supporting Information. The Fourier-transform (FT) magnitudes of the Ni K-edge EXAFS spectra for the charged LixNi0.8Co0.15Al0.05O2 with x = 0.5 and 0.1, as a function of temperature, are shown in Figure 8b. A NiO reference spectrum is also included for comparison. The strong first and second shells, located at ∼1.5 and 2.5 Å, respectively, are attributed to the single scattering path of the closest oxygen anions (i.e., Ni−O) and the second neighboring TM cations (i.e., Ni-M, where M = Ni, Co, Al) surrounding the absorbing Ni atom. The FT features of the EXAFS spectra for both samples indicate the increased formation of the NiO rocksalt structure around the Ni cations with increasing temperature. This is seen in the evolution of FT peaks corresponding to the NiO structure in the R range from 3 Å to 6 Å (Figure 8b). This is further verified from the bond length variations of the first Ni−O shell and the coordination numbers for the second Ni−M shell with temperature, shown in Figure 8c (subpanels (i) and (v)). As the temperature increases, the Ni− O bond lengths (∼1.87 Å, Ni4+−O) for both charged samples expand close to 2.07 Å, a value which corresponds to the Ni2+− O bond in the rock-salt NiO structure. The bond length for the second Ni−M shell (shown in Figure 8c, subpanel (ii)) also increases with temperature in a similar manner. This suggests that a bond length expansion occurs around Ni atoms due to the reduction of Ni ions from 4+ to 2+. As a result, the coordination number for the second Ni−M shell is increased from ∼6 for the layered (i.e., ordered rock-salt) structure to ∼12 for the NaCl-type rock-salt structure. The reduction of the Ni oxidation state and corresponding local structure changes (e.g., increase in the Ni−O bond length

and Ni−M coordination number) due to the thermal decomposition are initiated at significantly lower temperatures in the highly charged Li0.1Ni0.8Co0.15Al0.05O2 in comparison with the less-charged Li0.5Ni0.8Co0.15Al0.05O2. For instance, the Ni−O bond length for the Li0.1Ni0.8Co0.15Al0.05O2 increased by ∼0.15 Å (from 1.87 Å to 2.03 Å) from 150 °C to 300 °C, whereas the same bond length for the Li0.5Ni0.8Co0.15Al0.05O2 showed a very small expansion of 0.03 Å (from 1.88 Å to 1.91 Å) in the same temperature range. This is consistent with the TR-XRD/MS and Ni K-edge XANES results. In case of the highly charged Li0.1Ni0.8Co0.15Al0.05O2, noticeable peak broadening is observed in the first two FT peaks (Figure 8b) in this temperature range (150−300 °C). A corresponding increase of the Debye−Waller factors (Figure 8c, subpanels (iii) and (iv)) in this temperature range strongly suggests increased disorder in the local structure around Ni due to the massive oxygen loss in this region (see prior TR-XRD/MS results in Figure 4). Figure 9a presents the FT magnitude plots from the Co Kedge EXAFS spectra for the charged LixNi0.8Co0.15Al0.05O2, with x = 0.5 and 0.1, as a function of temperature. Shown for comparison are reference CoO and Co3O4 spectra. Compared with the Ni K-edge EXAFS results in Figure 8, the FT feature differences of the Co K-edge EXAFS spectra between the lesscharged Li 0.5 Ni 0.8 Co 0.15 Al 0.05 O 2 and highly charged Li0.1Ni0.8Co0.15Al0.05O2 samples are much more significant. In contrast to the formation of the NiO rock-salt structure around Ni in the less-charged Li0.5Ni0.8Co0.15Al0.05O2 at 450 °C (Figure 8a), the local structure around Co in the same sample did not fully transform to the CoO rock-salt structure at 450 °C (Figure 9a). This is more clearly confirmed by the variation of the Co−O bond length with temperature, presented in Figure 9b, showing a very small Co−O bond length expansion of 347



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Figure 10. In situ TEM images (a) before heating (25 °C) and (b) after heating up to 400 °C for a Li0.5Ni0.8Co0.15Al0.05O2 particle; in situ TEM images (c) before heating (25 °C), after heating at (d) 200 °C, and (e) 400 °C for a Li0.1Ni0.8Co0.15Al0.05O2 particle.

expansion of the first Co−O bond length, which is close to the value for CoO above 350 °C (Figure 9b, subpanel (i)). However, the temperature associated with the formation of the rock-salt structure local to Co is higher than for Ni, as is apparent in the Ni and Co K-edge FT spectra (see Figures 8b and 9a) and their fitting results (Figures 8c and 9b). This demonstrates that the Co ions are more thermally stable than the Ni ions in the charged LixNi0.8Co0.15Al0.05O2, which is in good agreement with the Ni and Co K-edge XANES results. It is noteworthy that the FT peak broadening and appearance of the shoulder peak at ∼2.9 Å for the second Co−M shell in the Li0.1Ni0.8Co0.15Al0.05O2 were quite prominent at 250 and 300 °C. Compared with the FT peaks of the reference Co3O4, these features are associated with the tetrahedral coordination of Co. This suggests formation of the Co3O4 spinel-like local structure around Co in this highly charged sample at 250 and 300 °C. Compared with the Ni K-edge fitting results, the relatively larger Debye−Waller factors for the Co−M shell (Figure 9b, subpanel (iv)) and large error bars in the coordination numbers for the Co−M shell (Figure 9b, subpanel (v)) at these temperatures (250 and 300 °C) indirectly suggest that the formation of a different local structure (e.g., Co3O4 spinel) around Co, which is not a layered or rock-salt structure. In other words, because only a rock-salt model was considered in the fitting, larger error bars suggest the formation of a nonrocksalt phase. Based on the Ni and Co K-edge XANES and EXAFS results, Ni prefers to directly form a rock-salt structure, whereas Co tends to stabilize by forming a Co3O4 spinel-like structure before undergoing a further phase transition to the rock-salt structure upon increasing temperature. In Situ Transmission Electron Microscopy (TEM). In situ TEM was used to study the nanoscale morphological changes occurring in the LixNi0.8Co0.15Al0.05O2, with x = 0.5 and 0.1, as a function of temperature. Figure 10a presents a brightfield micrograph of one of the Li0.5Ni0.8Co0.15Al0.05O2 particles

∼0.06 Å (from 1.88 Å to 1.94 Å) in the Li0.5Ni0.8Co0.15Al0.05O2 after heating to 450 °C. This is much less than the noticeable expansion in the Ni−O bond length of ∼0.16 Å (from 1.89 Å to 2.05 Å) in the same sample. This indicates that the local structure around Co in the Li0.5Ni0.8Co0.15Al0.05O2 may tend to be stabilized in a structure close to the Co3O4 spinel with less reduction of Co after heating to 450 °C, if the majority of the TM cations are Co. However, the FT peaks for the Li0.5Ni0.8Co0.15Al0.05O2 at 450 °C are not close to those of the Co3O4 reference, especially for peaks beyond the second Co− M shell. These FT peaks more or less resemble the CoO rocksalt structure (marked as “+” in Figure 9a). This is the main reason we refer to this as a Co3O4 spinel-like structure. It only shows a tendency toward forming the Co3O4 structure by the migration of some Co cations to the tetrahedral sites, but does not form the stoichiometric Co3O4 spinel structure beyond the first Co−O shell. This is likely due to the overwhelming contribution of the NiO rock-salt structure formed around Co in the high R range from 2 Å to 6 Å. This is dominant because the Ni content is ∼5 times larger than the Co content in the LixNi0.8Co0.15Al0.05O2. These results suggest that Ni and Co cations do not segregate to form local clusters, but are completely mixed together during the course of the thermal decomposition. This assumption is further supported by the nonsynchronized bond length variations of the first Co−O shell (Figure 9b, subpanel (i)) and the second Co−M shell (Figure 9b, subpanel (ii)) for the Li0.5Ni0.8Co0.15Al0.05O2 with increasing temperature. The bond length of the second Co−M shell (which contains contributions from the Ni reduction) expanded nearly twice as much as the first Co−O bond expansion (which contains contributions from Co reduction only) in the Li0.5Ni0.8Co0.15Al0.05O2. For the highly charged Li0.1Ni0.8Co0.15Al0.05O2, the formation of the CoO rock-salt structure was clearly observed in the FT features, marked as “+” in Figure 9a. It is also noticeable in the 348



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illustration depicting the phase stability map of the charged LixNi0.8Co0.15Al0.05O2 with x = 0.5, 0.33, and 0.1 during heating, based on the TR-XRD results in Figures 2−4. The starting temperature (the line marked as “①”) of the phase transition from the layered to the disordered spinel phase is quite flat, indicating a weak influence of the Li content (i.e., state of charge). In contrast, for the completion of this first phase transition (the line marked as “②”), the corresponding temperatures for more highly charged samples with x = 0.33 and 0.1 are much lower than that for x = 0.5 sample. The observed differences in the phase transition behaviors appear to originate from the activation barrier of cation migration occurring in the LixNi0.8Co0.15Al0.05O2 during heating, as described in the following section. Figure 12 shows the possible TM cation migration paths in the charged LixNi0.8Co0.15Al0.05O2 during the first phase transition from the layered to the disordered spinel phase. In the initial layered structure (Figure 12a), the TM cations occupy octahedral sites (TMoct layer) and the Li+ ions occupy the alternate layer of octahedral sites (Lioct layer).30 To complete the phase transition from the layered to the disordered spinel, some of the TM cations (most likely, Ni ions from XANES results) need to migrate from their original sites (labeled “A” in Figure 12b) to the octahedral sites in the Li layer (B in Figure 12(b)). This must be accompanied by the displacement of Li+ ions from their original octahedral sites to the adjacent tetrahedral sites, as shown in Figure 12c. The migration of TM cations can take two different paths, as shown in Figure 12d. One path is through a nearest neighboring tetrahedral site via the faces that it shares with the neighboring octahedra (path 1). The other path is a direct path from one octahedral site to the nearest neighbor octahedral site (path 2). Because the activation barrier for TM cation migration through the tetrahedral site in path 1 is reported to be smaller than the barrier of direct path 2,12,45 we can assume that most TM cation migration will take place through path 1. Since the tetrahedral sites of initial layered structure of the charged LixNi0.8Co0.15Al0.05O2 are empty, the activation energy for migration of TM cations from the octahedral site to the tetrahedral site (path 1-a) should be nearly the same for all

at room temperature. The particle has a dense morphology and is several hundred nanometers in diameter. The overall shape and morphology of the Li0.5Ni0.8Co0.15Al0.05O2 did not undergo significant changes, even after heating to 400 °C (Figure 10b). Figures 10c−e present bright-field TEM images of a more highly charged Li0.1Ni0.8Co0.15Al0.05O2 particle during heating from room temperature (Figure 10c) to 400 °C (Figure 10e). A porous structure with nanometer-sized holes in both in the bulk and at the surface was formed at ∼200 °C (Figure 10d) and sustained up to 400 °C (Figure 10e). This is in direct contrast to the Li0.5Ni0.8Co0.15Al0.05O2 particle, which maintained its initial, dense structure up to 400 °C and only exhibited slight changes in morphology at the near surface region. The TRXRD/MS results for the Li0.1Ni0.8Co0.15Al0.05O2 samples (Figures 4a−c) showed a substantial release of both oxygen and CO2, starting at ∼200 °C. It can be concluded that the significant, rapid loss of oxygen from the Li0.1Ni0.8Co0.15Al0.05O2 material is correlated with the formation of the porous structure during thermal treatment. Phase Stability and Cation Migration Mechanism during Phase Transitions. Figure 11 is a schematic

Figure 11. Scheme of the phase transitions of LixNi0.8Co0.15Al0.05O2 during thermal decomposition based on the TR-XRD data in Figures 2−4.

Figure 12. Phase transition of the LixNi0.8Co0.15Al0.05O2 charged cathode during heating: (a) layered structure, (b) cation migration during the phase transition from the layered structure to the disordered-spinel structure, (c) disordered spinel structure, and, (d) cation migration path from octahedral A to octahedral B. 349



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and Co K-edge XANES and EXAFS results, as a function of temperature, showed that the release of oxygen is closely related to the migration of Ni and Co cations in the cathode materials, and these migrations affect the kinetics of the phase transition from the layered to the disordered spinel and, subsequently, to the rock-salt structure. In addition, it was found that Ni prefers to form a rock-salt structure directly, whereas Co tends to stabilize by forming a Co3O4 spinel-like structure, before a further phase transition to a complete rocksalt structure upon further heating. In situ TEM during heating has showed the formation of a highly porous structure of the highly overcharged Li0.1Ni0.8Co0.15Al0.05O2 sample caused by the rapid loss of oxygen during thermal decomposition. This systematic study, using a combination of advanced diagnostic tools, provides new insight into the thermal decomposition mechanism of charged cathode materials, including direct correlations between phase transitions and O2 and CO2 gasreleasing behavior, cation migration pathways, and microscopic morphology changes. The information gained in this study will guide engineers and scientists to the rational design of thermally stable cathode materials for practical lithium-ion battery systems.

three samples. The slight decrease in the observed phase transition temperature that occurs with less lithium content might be due to the differences in the valence state of the TM cations (e.g., Ni3+ vs Ni4+), which should have a major effect on the ionic mobility in a cubic close-packed oxygen framework.12 Therefore, the onset temperatures of the phase transition from layered to disordered spinel phase for LixNi0.8Co0.15Al0.05O2 are close to each other. However, the second part of the TM cation migration from the tetrahedral site to the octahedral site (path 1-b) is strongly dependent on the Li content, because of the differences in the occupancy of Li+ ions in the octahedral sites. For x = 0.33 and x = 0.1 samples (with less Li+-ion occupancy in the Lioct layer), the TM cations can migrate easily from tetrahedral sites to octahedral sites, resulting in a more-rapid completion of the phase transformation from the layered to the disordered spinel phase. In addition, the large oxygen release from the Li0.1Ni0.8Co0.15Al0.05O2 is likely responsible for the rapid completion of the phase transition from the layered to the disordered spinel phase. The large amount of released oxygen would create a significant number of oxygen vacancies, thereby lowering the activation barrier for the migration of TM cations from the tetrahedral sites to the octahedral sites in the Lioct layer. This effect would be further enhanced by the large number of vacant Li sites created by the high degree of delithiation. After the formation of the disordered spinel structure, the phase transition to the NiO-type rock-salt structure takes place upon additional oxygen release. Through this phase transition, the Li and TM cations are all randomly distributed at the octahedral sites. The evolution of (220)s peak intensity in the TR-XRD data shows that the TM cations at tetrahedral sites immediately moved to octahedral sites during the phase transition from the disordered spinel to the rock-salt phase. Most of the cations occupying the tetrahedral sites are Co, according to our interpretation of the Co K pre-edge XANES and EXAFS analysis. The M3O4-type spinel, with some Co cations occupying the tetrahedral sites, could push the phase transition to the rock-salt structure to a higher temperature.48

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ASSOCIATED CONTENT

S Supporting Information *

Constant current charge profiles, Rietveld refinement of XRD, MS results for O2 release for the LixNi0.8Co0.15Al0.05O2 (x = 0.5, 0.33 and 0.1), and detailed description of EXAFS fitting analysis. This material is available free of charge via the Internet at http://pubs.acs.org.

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AUTHOR INFORMATION

Corresponding Author

*E-mails: [email protected] (K.N.), [email protected] (K.C.), [email protected] (X.Y.). Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS The work done at Brookhaven National Lab. was supported by the Assistant Secretary for Energy Efficiency and Renewable Energy, Office of Vehicle Technologies of the U.S. Department of Energy (DOE), under Contract No. DE-AC02-98CH10886. The work done at KIST was supported by Global Research Laboratory Program, through the National Research Foundation of Korea (NRF), funded by the Ministry of Education, Science and Technology (MEST) (Grant No. 2010-00351). The work done at Yonsei University was supported by the Converging Research Center Program through the Ministry of Education, Science and Technology (No. 2012K001266). Electron microscopy was performed at the Center for Functional Nanomaterials (CFN); XRD and XAS were carried out at the Nationsl Synchrotron Light Source (NSLS), Brookhaven National Laboratory. The CFN and NSLS are supported by the U.S. Department of Energy, Office of Basic Energy Sciences, under Contract No. DE-AC02-98CH10886.

CONCLUSION We have developed a new diagnostic technique that uses a combination of in situ time-resolved X-ray diffraction (TRXRD) and mass spectroscopy (MS) to study the thermal decomposition mechanism of charged lithium battery cathode materials. This new technique has allowed the direct correlation between structural changes and gas evolution during the thermal decomposition of LixNi0.8Co0.15Al0.05O2 cathode materials charged to x = 0.5, 0.33, and 0.1 levels. This study has shown that the state of charge affects both the structural changes that occur in these materials, as well as the evolution of O2 and CO2 gases during thermal decomposition. The evolution of both O2 and CO2 gases are well-correlated with the phase transitions that occur during thermal decomposition. Overcharging (i.e., overdelithiation) causes higher thermal instability and larger amounts of oxygen release at lower temperatures. In particular, highly overcharged cathode samples (i.e., Li0.1Ni0.8Co0.15Al0.05O2) show a severe oxygen release at ca. 175 °C, accompanied by a sudden structural change from the layered to the disordered spinel phase. Because the released oxygen is highly reactive, it leads to the decomposition of carbon either from the PVDF binder or the conducting carbon in the charged electrodes, resulting in additional CO 2 production at much lower temperatures than anticipated. Ni

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REFERENCES

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