Facilitating the Operation of Lithium-ion Cells with High-nickel Layered

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Facilitating the Operation of Lithium-ion Cells with Highnickel Layered Oxide Cathodes with a Small Dose of Aluminum Jianyu Li, Wangda Li, Shanyu Wang, Karalee Jarvis, Jihui Yang, and Arumugam Manthiram Chem. Mater., Just Accepted Manuscript • DOI: 10.1021/acs.chemmater.8b01077 • Publication Date (Web): 20 Apr 2018 Downloaded from http://pubs.acs.org on April 20, 2018

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Chemistry of Materials

Facilitating the Operation of Lithium-ion Cells with High-nickel Layered Oxide Cathodes with a Small Dose of Aluminum Jianyu Li,a Wangda Li,a Shanyu Wang,b Karalee Jarvis,a Jihui Yang,b Arumugam Manthirama * a

McKetta Department of Chemical Engineering & Texas Materials Institute, The University of Texas at Austin, Austin, Texas, 78712, USA b Material Science and Engineering Department, University of Washington, Seattle, Washington, 98105, USA

Abstract Layered oxide cathodes with a high Ni content of > 0.6 are promising for high-energy-density lithium-ion batteries. However, parasitic electrolyte oxidation of the charged cathode and mechanical degradation arising from phase transitions significantly deteriorate the cell performance and cycle life as the Ni content increases. We demonstrate here a significantly prolonged cycle life with superior cell performance by substituting a small-dose of Al (2 mol.%) for Ni in LiNi0.92Co0.06Al0.02O2; the capacity retention after operating a full cell fabricated with graphite anode for 1,000 cycles increases from 47% to 83% on going from the Al-free LiNi0.94Co0.06O2 to the Al-doped LiNi0.92Co0.06Al0.02O2 cathode. Through in-situ X-ray diffraction, we provide the operando evidence that the Al-doping tunes the H2-H3 phase transition process from a two-phase-reaction to a quasi-mono-phase reaction, minimizing the mechanical degradation. Furthermore, secondary-ion mass spectrometry reveals considerably suppressed transition-metal dissolution with Al-doping, effectively preventing sustained parasitic reactions and active Li trapping due to chemical crossover on graphite anodes. This work offers a viable approach for adopting high-Ni cathodes in lithium-ion batteries.

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1. Introduction With a dramatically increasing demand for electric vehicles (EVs), a driving range of minimum 300 miles is a must to satisfy the market requirement.1 Such a demand has motivated enormous efforts focused on electrode materials, in particular cathode materials for lithium-ion (Li-ion) batteries.2-5 However, those cathode materials that can be practically applied in the automotive industry in the next 10 years are more likely limited to a small number of compounds.1 Among them, high-Ni layered oxide LiNi1-xMxO2 (M=Mn and Co) with (1-x) > 0.6 is one of the most promising candidates. Generally, as the Ni content in LiNi1-xMxO2 increases, Li+ diffusivity, electronic conductivity, rate capability, and discharge capacity increase.6 Nevertheless, the disadvantages will become more challenging. Consider LiNiO2 for example, formation of stoichiometric LiNiO2 with a Li/Ni ratio of 1 : 1 is difficult as Ni3+ tends to get reduced to Ni2+ at the synthesis temperature, which leads to Li/Ni cation mixing in the lithium layer to from [Li1-yNiy]3a[Ni]3b[O2]6c due to the similar sizes of Li+ and Ni2+.7 The anti-site Ni2+ in the Li layer greatly hinders Li-ion transport, resulting in poor electrochemical performance. In addition, during electrochemical cycling, the generation of monoclinic and hexagonal structures in the delithiated state8 induces large unit cell volume changes and anisotropic strains, which lead to micro-crack formation and eventually secondary particle disintegration.9 This structural evolution process could be more severe as Ni content increases due to the presence of hexagonal 3 (H3) phase in deeply delithiated state.10 Meanwhile, a high concentration of unstable Ni4+ in the highly delithiated LiyNiO2 tends to react aggressively with the carbonate electrolyte.11 The unwanted reaction greatly expedites the irreversible transformation of the layered structure to a rock-salt (NiO) phase. Overall, these various degradation processes result in a high interfacial impedance and poor electrochemical


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performance.7，12-14 To tackle these issues, Co and Al co-substitution into the Ni site, which gives rise to the NCA (LiNi1-y-zCoyAlzO2) system, has been proven an effective strategy.15-17 Most studies on the NCA system have been focused on materials with Ni content less than 0.8.18-24 However, increasing Ni content is preferred to increase the energy density of Li-ion batteries. On the other hand, Al content was conventionally set to be no less than 0.05 in the transition-metal site to stabilize the structure.25-27 However, Al3+ is electrochemically inactive and it reduces the capacity and energy density. We present here a new NCA (LiNi0.92Co0.06Al0.02O2, denoted as NNCA) with a high Ni content of 0.92 and a low Al content of 0.02. Compared with the baseline material – LiNi0.94Co0.06O2 (denoted as NC) and LiNi0.8Co0.15Al0.05O2 (denoted as CNCA), NNCA exhibits superior electrochemical performance with a discharge capacity of as high as 217 mA h g-1 and an excellent capacity retention of 83% even after 1000 cycles in pouch-type full cells. NNCA also exhibits outstanding rate capability, delivering a superb discharge capacity of 123 mA h g-1 at 30C. Various characterizations including ex- and in-situ X-ray diffraction (XRD), X-ray photoelectron spectroscopy (XPS), high resolution transmission electron microscopy (HRTEM), diffraction scanning transmission electron microscopy (D-STEM), and time-offlight secondary-ion mass spectrometry (TOF-SIMS) have been carried out to investigate the origin of the markedly enhanced electrochemical performance. It is proved that such a small amount of Al in NNCA plays a critical role in stabilizing the cathode-electrolyte surface and mitigating the detrimental structure evolutions. 2. Experimental Procedures 2.1 Precursor preparation Ni0.94Co0.06(OH)2 precursor with a particle size of ≈ 12µm was prepared by a transitionmetal co-precipitation route. Typically, NiSO4·6H2O and CoSO4·7H2O with a molar ratio of 94 : 6 (Ni : Co) were dissolved in distilled water to form a combined concentration of 1.0 M. 3 ACS Paragon Plus Environment


Then the mixed solution was pumped into a 2.5 L continuous stirring tank reactor (CSTR), with separate feeds of solution containing a proper amount of saturated NH4OH and KOH. The PH (~ 11) and temperature (~ 50℃) were carefully controlled during the whole reaction. Finally, the co-precipitated powder was filtered, washed with distillated water several times, and then dried at 100℃ overnight. Al(OH)3-coated Ni0.94Co0.06(OH)2 was prepared by a solution-based process. A stoichiometric amount of Al isopropoxide was dissolved in a proper amount of isopropanol at 60℃. Then the solution was mixed vigorously with Ni0.94Co0.06(OH)2 precursor and stirred to dry. The resulting powder was dried at 100℃ overnight and then sieved for future use. 2.2 LiNi0.94Co0.06O2 (NC) and LiNi0.92Co0.06Al0.02O2 (NNCA) Powder Preparation The obtained Ni0.94Co0.06(OH)2 precursor was mixed with LiOH·H2O (Li : Ni + Co = 1 : 1.05 by mole), followed by calcinating at 500℃ for 5 h and then heating at 640℃ for 15 h with flowing oxygen to form NC. To produce NNCA, the Al(OH)3-coated Ni0.94Co0.06(OH)2 precursor was also uniformly mixed with LiOH·H2O (Li : Ni + Co + Al = 1 : 1.05 by mole). The mixture was then calcinated at 500℃ for 5 h, followed by heating at 680℃ to yield NNCA. 2.3 LiNi0.8Co0.15Al0.05O2 (CNCA) Powder Preparation Ni0.85Co0.15(OH)2 precursor with a particle size of ≈ 12µm was prepared by the same transition-metal co-precipitation route as that for Ni0.94Co0.06(OH)2 except that NiSO4·6H2O and CoSO4·7H2O were mixed with a molar ratio of 85 : 15 (Ni : Co). Then Ni0.85Co0.15(OH)2 powder was coated with a proper amount of Al(OH)3 by a solution-based process. The obtained Al(OH)3-coated Ni0.85Co0.15(OH)2 was mixed with LiOH·H2O (Li : Ni + Co = 1 : 1.03 by mole), followed by calcinating at 750℃ for 15 h with flowing oxygen to form CNCA.


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2.4 Material Characterization Powder X-ray diffraction (XRD) measurement was carried out with a Miniflex 600 X-ray diffractor with Cu Kα radiation at a scan rate of 0.38°/min from 10° to 80°. The chemical compositions were determined by inductively coupled plasma optical emission spectrometer (ICP-OES, Varian 715 ES). Morphology was examined with a field-emission scanning electron microscope (SEM, Quanta 650). A JEOL 2010F TEM was used to conduct elemental mapping with energy disperse X-ray spectroscopy (EDS) and structural analysis by high resolution TEM (HRTEM) and D-STEM with a 1 − 2 nm probe. The TEM sample was prepared by thinning a spherical particle to a thickness less than 100 nm with a FEI Strata™ DB235 dual-beam SEM/FIB system. X-ray photoelectron spectroscopy (XPS) measurement was performed with an Axis Ultra DLD spectrometer (Kratos) using Al Kα radiation. In-situ X-ray diffraction study was carried out using a Bruker D8 Advance X-ray diffractometer equipped with a Cu Target X-ray tube and a LYNXEYE XE high-resolution energydispersive 1D detector. The in-situ cells were cycled between 2.8 V and 4.4 V at ~ 25oC using a Maccor Series 4200 M 16-Channel Automated Battery Test System (Maccor, USA). Diffraction patterns were collected in the scattering angle (2θ) ranges of 18 − 21°, 36 − 40°, 43 − 52°, and 57 − 72° at 0.02° intervals with a dwell time of 0.6 s. Each scan took ~ 16 min. The diffraction peaks of Be, BeO, Al, PVDF, and carbon black may exist in the diffraction patterns, yet these peaks do not shift during the in-situ test. TOF-SIMS was carried out by a TOF.SIMS 5 spectrometer (ION-TOF GmbH). All detected secondary ions of interest had a mass resolution of >5000 and possessed negative polarity. A pulsed 30 keV Bi1+ (20 ns) ion beam set in high current mode was applied for depth profiling and a 500 eV Cs + (negative) ion beam was used for the sputtering of the cycled electrodes with a typical sputtered area of 300 µm

300 µm. The typical analyzed area was 100 µm


100 µm.


2.5 Electrode Preparation and Electrochemical Characterization The cathode was prepared as follows: the active material (NC, CNCA, or NNCA), Super P conductive carbon, and poly(vinylidene) fluoride (PVDF) with the weight ratio of 8 : 1 : 1 were dispersed in N-methyl-2-pyrolidone (NMP) and stirred to form uniform slurry. The obtained slurry was cast on an Al foil with an active-material loading of ~ 4 mg cm-2, followed by drying and cutting into circular disks. CR2032 coin cells were assembled inside an Ar-filled glove box. Li metal was utilized as anode for the half-cell test. The electrolyte solution is 1.0 M LiPF6 in a mixture of ethylene carbonate/ethyl methyl carbonate (EC : EMC = 3 : 7 by weight) with 2 wt.% vinylene carbonate (VC). NC and NNCA half cells were cycled at a constant rate of C/5 between 4.4 V and 2.7 V. The voltage window is between 4.5 V and 2.5 V for CNCA half cells. Long-term cycling test was performed with laminatedpouch-type full-cells consisting of NC or NNCA as the cathode and mesocarbon microbead (MCMB) graphite as the anode. The N/P ratio was controlled to be ~ 1.1. Two formation cycles were performed at C/10, followed by cycling at C/2 between 4.3 V and 2.7 V for full cells. 3. Result and Discussion 3.1 Al Distribution and Structure Characterization of Pristine Material Figure 1a shows the SEM images of the as-prepared NC and NNCA. Both samples possess spherical morphologies with a secondary-particle size of 12 µm in average. The chemical compositions of the as-prepared NC and NNCA were confirmed by inductively-coupled plasma atomic emission spectroscopy (ICP-AES) (Table S1, Supporting Information). The chemical formula is Li0.9Ni0.94Co0.06O2-δ for NC and Li0.95Ni0.92Co0.06Al0.02O2-δ for NNCA. The bulk structure of NC and NNCA was characterized by ex-situ XRD, as shown in Figure 2a. XRD spectra of both NC and NNCA can be indexed to a α-NaFeO2 layered structure with a R 3 m space group and no impurity is detected.17 Also, NNCA exhibits a higher I(003)/I(104) 6 ACS Paragon Plus Environment

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ratio (1.28 for NNCA and 1.20 for NC), indicating decreased cation.28,29 In addition, the peak splitting of (006)/(102) and (108)/(110) pairs in NNCA is more distinguishable, suggesting a more ordered layered structure. Figure 1b and 1c depict the STEM images and corresponding energy disperse X-ray spectroscopy (EDS) maps of NNCA. Elements including Ni, Co, and Al are all uniformly distributed throughout the particle without elemental segregation. This is further confirmed by the EDS spot scans on the secondary particles from the surface towards the bulk as shown in Figure S15. The content of Al is consistent from spot to spot, indicating that there is no concentration gradient. Also, comparative intensity distribution of AlO-, 62

NiO-, and

58

NiO- fragments on a fresh NNCA surface is depicted in Figure S14, and the

result shows the surface is dominated by

62

NiO-/58NiO- fragments instead of AlO-. The

normalized depth profile (the inset in Figure S15) shows good overlap of AlO-, 58

62

NiO-, and

NiO- fragments, suggesting that the distributions of the fragments is uniform throughout the

particles. Additionally, D-STEM was employed to check the structure and it reveals a R 3 m phase throughout the primary particle of NNCA, including the surface (Figure 2b). Thus, the combined TOF-SIMS and D-STEM analysis shows that Al does not segregate on the surface. According to the above analyses, it is reasonable to conclude that the doped Al forms a solid solution with Ni and Co in the transition-metal site and is homogeneously distributed within the NNCA particles. 3.2 Electrochemical Performance The superior electrochemical performance of NNCA compared to those of NC and CNCA is first illustrated by their typical discharge curves, as shown in Figure 3a. Within a wider voltage window (2.5 V – 4.5 V vs. Li/Li+), CNCA can only deliver a discharge capacity of 200 mA h g-1, which is 17 mA h g-1 lower than NNCA and 26 mA h g-1 lower than NC. Moreover, Figure 3b displays the discharge capacities of the three samples at different discharge rates. Initially, all three samples demonstrate similar rate capability from C/10 to 5C, 7 ACS Paragon Plus Environment


with a lower discharge capacity for CNCA. However, the rate performance diverges at higher C rates and NNCA exhibits much better rate capability than NC and CNCA. At 10C rate, NNCA approaches a high discharge capacity of 172 mA h g-1 (78% capacity retention based on the discharge capacity at C/10), considerably outperforming NC (160 mA h g-1, 69% retention) and CNCA (124 mA h g-1, 59% retention). Even at 30C, NNCA still attains a discharge capacity of as high as 123 mA h g-1 as compared to NC (99 mA h g-1) and CNCA (no detectable capacity). The improved rate capability of NNCA is not only a result of high Ni-content (low Li-ion transport activation barrier around Ni cations), 30 but also attributed to a more ordered structure in NNCA. Figure 3c depicts the cycling performance of NC, NNCA, and CNCA in half cells. The discharge capacities of NC and NNCA rise, respectively, from 206 to 226 mA h g-1 and from 202 to 217 mA h g-1 during the first 10 cycles. Both peak capacities are among the highest values reported in the high-Ni layered oxide system.31-33 Afterwards, NC suffers from much faster capacity decay. It retains a discharge capacity of 133 mA h g-1 after 200 cycles, corresponding to a low capacity retention of 58%. In sharp contrast, NNCA demonstrates a significantly improved cyclability, as evidenced by a high reversible discharge capacity of 195 mA h g-1 after 200 cycles, corresponding to a capacity retention of as high as 90%. The superior cyclability of NNCA is also reflected in the high Coulombic efficiency (CE) (Figure S2, Supporting Information). As for CNCA, it delivers slightly lower capacity retention than NNCA – 87% capacity retention after 200 cycles with a 0.1 V higher operating voltage. Figure 3d and 3e plot the charge/discharge profiles of both NC and NNCA from the 10th to the 200th cycle, respectively. At the 10th cycle, both samples exhibit similar voltage curves with distinguishable plateaus around 4.0 and 4.1 V, corresponding to the hexagonal 1 to hexagonal 2 (H1-H2) and hexagonal 2 to hexagonal 3 (H2-H3) phase transition process, respectively. After 200 cycles, however, both the previously observed plateaus disappear in 8 ACS Paragon Plus Environment

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NC, implying a severely degraded crystal structure. In contrast, NNCA still displays wellpreserved plateaus. Furthermore, NC shows a rapid decrease in the average discharge voltage after 200 cycles (declined from 3.83 V to 3.63 V, dashed line in Figure 3d), while the average discharge voltage of NNCA only drops by 0.06 V (dashed line in Figure 3e). Such stabilized voltage behavior in NNCA contributes to a high energy-density retention of 89% after 200 cycles, greatly outperforming that of NC (55% retention, Figure S3, Supporting Information). It is worth noting that current Li-ion batteries deliver an energy density of ~ 250 Wh kg-1 at the cell level. 34 But an energy density of 300 Wh kg-1 at the cell level is needed to power an EV for 300 miles, and this requires an energy density of 750 Wh kg-1 at cathode-material level (around 202 mA h g-1 at 3.7 V vs. Li/Li+).1，11 Our NNCA is capable of achieving an energy density over 800 Wh kg-1 with outstanding capability retention, fulfilling the high energy requirement of EVs. In consideration of practical application, NC and NNCA were paired with the graphite anode in pouch-type full cells for long-term cycling test. As shown in Figure 3f, NNCA delivers a high capacity retention of 83% even after 1,000 cycles. In comparison, NC only retains 47% of its original capacity. As for voltage retention, NNCA has a voltage drop of only 0.17 V with a high voltage retention of 95% after 1000 cycles, while NC undergoes a much larger voltage drop of 0.57 V (Figure S4, Supporting Information). Such a sharp contrast clarifies that excellent energy-density retention can also be achieved in full cells employing NNCA. 3.3 Surface Reactivity Analysis To investigate the origin of the significantly improved electrochemical performance of NNCA, self-discharge experiments were carried out to evaluate the surface stability. The selfdischarge of high-Ni electrode undergoes the following reaction: 35



Li x MO 2  y Li   y Electrolyte  Li x  y MO 2 + y Electrolyte + (M = Ni, Co, and Al)

In this work, Li│NC or Li│NNCA half cells were charged to 4.4 V and then rested for 12 h at 55℃, followed by discharging to 2.8 V. During the charge process (Figure 4a), both samples deliver abnormally high charge capacities, mainly resulting from the aggravated side reaction between the cathode and electrolyte at elevated temperature. Figure 4b illustrates the variations of open circuit voltage (OCV) during storage time. NC suffers from a rapid OCV drop and ends up with 3.48 V, implying a fast relithiation process. After storage, it only obtains a discharge capacity of < 10 mA h g-1 (Figure 4c). However, NNCA exhibits much more stable behavior. Its OCV could be stabilized at 3.91 V after 12 h of storage, and it retains a high discharge capacity of 120 mA h g-1 after storage. The self-discharge test evidently demonstrates that Al doping in NNCA is successful in alleviating the reactivity of high-Ni layered oxide cathode. Furthermore, XPS and TOF-SIMS were conducted to inspect the differences in the surface chemistry of the cycled cathodes and graphite anodes. The elemental concentration profile of the cycled cathodes is summarized in Figure 5a. The concentration of O element increases after Al doping, which could arise from more O-M bonding (Figure 5b). Additionally, the Ni content in NNCA is over twice higher than NC (4.7 at.% for NNCA and 2.2 at.% for NC). The increased O-M bonding content and Ni content together evidence a thinner cathodeelectrolyte interphase (CEI) for NNCA, 36-38 indicative of a less reactive surface that benefits cycling performance. More importantly, the concentration of F element declines from 18.9 to 7.2 at.% after Al doping, which is mainly due to the less LiF/MFx resulting from dissolved transition-metal cations(Figure 5c). Note that the dissolved transition-metal species can migrate to the anode side, damaging the solid-electrolyte interphase (SEI) at the graphite anode,39 causing severe active Li trapping (including metallic lithium microstructures), and thus severely lowering the capacity retention.40 This is clearly revealed by TOF-SIMS 10 ACS Paragon Plus Environment

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analysis on graphite anodes after long-term cycling in Figure 6. A series of secondary-ion fragments of interest, including C2HO-, 7Li-,

58

Ni-, and Co- represent, respectively, the

electrolyte breakdown products, plated lithium (dead lithium), and transition metal dissolution species (both 58Ni- and Co-).41 As seen, all these fragments show significantly reduced signal yields for the NNCA-paired graphite, especially 7Li- (reduced by 90%), 58Ni- and Co- (reduced by, respectively, 85% and 75%). As a result, the small dose of Al in NNCA is remarkably effective in reducing active mass dissolution from the cathode, parasitic electrolyte reductive breakdown, and active lithium trapping on the graphite anode, which is crucial in prolonging service life. 3.4 Structural Evolution Analysis Aiming to better understand the phase transition process, in-situ XRD was performed to probe the detailed structural evolution of the cathode samples during cycling. Figure 7a depicts the time-resolved XRD patterns recorded during the first charge process. The contour plots the 2θ regions with Bragg reflections from (003) to (113). The structural behavior of both NC and NNCA upon delithiation is similar to previous studies42 – phase transitions from H1 to H2, and then to H3. Note that the shift in the (003) peak represents the lattice parameter change along the c-axis only and the shift in the (110) peak indicates the lattice parameter change along the a- and b-axes equally. Both samples exhibit similar distortion behaviors along the a- and b-axes. The (110) peak of NC and NNCA monotonically shifts to higher 2θ angles during the whole charge process, indicating a continuous shrinkage along a- and baxes due to the oxidation of Ni3+ to smaller Ni4+. The shrinkage becomes less intense as evidenced by the slight shift in the (110) peak at higher state-of-charge (SOC, actual capacity/first charge capacity). Finally, NC and NNCA reach similar (110) peak positions at full SOC (Figure S8, Supporting Information).



However, the lattice distortion behavior becomes more complicated along the c-axis as reflected in the (003) peak. During charging, the (003) peak of both samples initially shifts to lower 2θ angles, indicating the expansion of c-axis accompanied by the phase transition from H1 to H2. At high SOC, the (003) peak undergoes a much more rapid and aggressive shift to higher 2θ angles, suggesting a severe shrinkage along the c-axis with further phase transition to H3. As shown in Figure 7b, the (003) peaks of both NC and NNCA initially go through a smooth shift to a lower angle of 18.4o until a certain SOC (~ 70% SOC for NNCA and ~ 60% SOC for NC). Note that SOC is proportional to the time operated and both samples complete the charge process within the almost same period (10.06 h for NC and 10.01 h for NNCA), thus NC finishes the leftward shift in a shorter time, which induces more intense internal strains. Thereafter, the (003) peaks of both samples shift to higher 2θ angles. Interestingly, at an SOC higher than 80%, NC exhibits a two-phase reaction until the end of charge – a new (003)H3 peak emerges at higher 2θ angles while the original (003)H2 peak gradually disappears. In sharp contrast, NNCA goes through a quasi-single-phase reaction within the same region – the (003) peak of NNCA gradually shifts to higher 2θ values without any obvious appearance of a new peak. It is noted that at 83% SOC (Figure 7b), the 003 peak of NNCA is broad, which could possibly be due to the presence of two phases with close lattice parameters. However, the broadening of the other reflections is not apparent, suggesting that the twophase reaction is greatly undermined by aluminum doping. Also, the co-existence of two phases is accompanied by strong internal strains arising from the significant lattice mismatch between H2 and H3 phases. Accordingly, the internal strains tend to accelerate the structural rupture of NC during cycling. The quasi-single-phase reaction in NNCA greatly attenuate the strain and help maintain the structural integrity. Apart from the advantages of the H2-H3 phase transition mechanism discussed above, NNCA also exhibits superior structural reversibility. Figure 8a shows the contour plot of the 12 ACS Paragon Plus Environment

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(003) peak for two cycles. As indicated by the dashed line, the position of the (003) peak for NNCA remains almost unchanged at 4.4 V at the second cycle. However, the (003) peak of NC shifts to higher 2θ angles at 4.4 V at the second cycle. In addition, Figure 8b provides a more straightforward evidence. The (003)H3 peak separation between the two cycles in NC is as high as 0.36o. According to the Bragg Equation (2dsinθ = λ), this 0.36o difference stands for ~ 1.8% more shrinkage along the c-axis. In striking contrast, this separation is largely alleviated for NNCA – the (003)H3 peak at the second cycle is only 0.08o higher. Notably, at the second cycle, the (003)H3 peak in NC at 100% SOC is at 19.8o – 0.4o higher than that of NNCA (Figure S9a, Supporting Information). This 0.4o difference in NC indicates 2% more shrinkage along the c-axis than in NNCA, which is sufficient enough to deteriorate structural integrity and the electrochemical performance especially during prolonged cycling. Furthermore, as shown in Figure 8c, after two cycles, the (003) peak in NC apparently shifts by 0.08o to a lower 2θ value at 2.8 V, while this shift is almost indistinguishable in NNCA. Thus, with the incorporation of Al, NNCA exhibits significantly suppressed lattice distortion, less internal strain, and better structural reversibility. 4. Conclusion In summary, this work presents a new NCA (NNCA) cathode material with a high Ni content of 0.92 and highlights the possibility of stabilizing high-Ni layered oxides by doping only a small dose of Al (2 mol.% in the transition-metal layer). The as-prepared NNCA delivers a high discharge capacity of 217 mA h g-1, satisfying the requirement for the cathode material to reach to an energy density of 300 Wh kg-1 at the cell level. The practicality of NNCA is also demonstrated in pouch-type full cells paired with a graphite anode, delivering remarkable capacity retention (83%) and voltage retention (95%) after 1000 cycles. By virtue of Al incorporation, the surface chemical stability and bulk structural reversibility of NNCA are significantly enhanced. In-situ XRD provides evidence on the smoothened H2-H3 phase 13 ACS Paragon Plus Environment


transition of NNCA in reducing lattice distortions and accompanied internal strain. Moreover, inhibited transition-metal dissolution and migration to graphite anodes with Al incorporation in NNCA are revealed by TOF-SIMS, which minimizes parasitic reactions and active Li trapping under extensive full-cell operation. This work provides a viable approach for the development of high-Ni (Ni > 0.9) layered oxide cathodes with long cycle and calendar life for high energy-density Li-ion batteries. Associated Content Supporting Information. Rietveld refinement result and XRD patterns, additional electrochemical test results (Coulombic efficiency, specific energy, charge/discharge curves, discharge voltage, CV curves), SEM images, XPS data, in-situ XRD patterns, and ICP results.

Author Information Corresponding Author * E-mail: [email protected] Notes The authors declare no competing finacial interest. Acknowledgements This work was supported by the Assistant Secretary for Energy Efficiency and Renewable Energy, Office of Vehicle Technologies of the U.S. Department of Energy through the Advanced Battery Materials Research (BMR) Program (Battery500 Consortium) award number DE-EE0007762 and the Welch Foundation grant F-1254. The author also thanks Mr. Chang Da for assistance with in-situ XRD data visualization, Dr. Hugo Celio for assistance with XPS characterization and Dr. Andrei Dolocan for assistance with TOF-SIMS characterization.


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